EFFECTS OF AUSTEMTE C0M)ITION ON BAINITE TRANSFORMATIONS IN MICROALLOmD BAR
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EFFECTS OF AUSTEMTE C0M)ITION ON BAINITE TRANSFORMATIONS IN MICROALLOmD BAR STEELS BY Fen Zhang A thesis submitted to the Department of Materials and Metaiiurgical Engineering in confonnity with the requirements for the Degree of Doctor of Philosophy Queen's University Kingston, Ontario, Canada June 1997 copyright '%en Zhang, 1997
Transcript of EFFECTS OF AUSTEMTE C0M)ITION ON BAINITE TRANSFORMATIONS IN MICROALLOmD BAR
EFFECTS OF AUSTEMTE C0M)ITION ON BAINITE TRANSFORMATIONS IN
MICROALLOmD BAR STEELS
BY
Fen Zhang
A thesis submitted to the
Department of Materials and Metaiiurgical Engineering
in confonnity with the requirements for the
Degree of Doctor of Philosophy
Queen's University
Kingston, Ontario, Canada
June 1997
copyright '%en Zhang, 1997
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Effects of austenite condition on bainite transformations have been investigated for
2 rnicrodoyed bar steels: a O. 11C-0.003B-0.09M1 steel ('L') and a 0.44C-0.07%-0.03V
steel ('My). A quench-deformation dilatometer was used to subject the samples to a range
of austenite conditions produced by different thermomechanical processing. From
observations of the final microstructures and analyses to the diiation data, CCT diagrams
have been obtaiaed, bainite types classified, the effects of austenite condition on bainite
transformations clarified, and a bainite transformation mode1 pro posed.
For L steel, austenite condition has a significant effect on bainite transformation
kinetics and morphologies. Deformed and recrystaliized austenite generally results in a
delayed bainite reaction due to a decrease in bainite growth. Unrecrystallized austenite
leads to an accelerated bainite transformation due to increased nucleation sites and rate
and growth rate of bainite. For M steel, the effects of austenite condition on bainite
transformations are not as signlficant as in L steel due to the existence of grain-boundary
ferrite.
Dunng continuos cooling, bainite nucleates at austenite grain boundaries, twin
boundaries, pre-existing bainite laths, coarse precipitates, deformation bands and
dislocation subgrain boundaries. It is proposed that bainitic femte lengthens displacively,
and thickens by either a difision-controlled ledge growth at higher temperatures or a
coalescence of low-angle sublaths at lower temperatures.
ACKNOWLEDGMENTS
Firstly and most importantly, 1 would like to express my deep gratitude to
Professor I. Douglas Boyd for his strong support and guidance throughout the research
and thesis writing.
Secondly, 1 would like to thank Charlie Cooney, Paul Nolan, Darryl Dietrich and
Gary Contant for technical support; and Shirley Donnelly and Chns Fowler for constant
support. Marek Marchwica in Stelco and Bill Heitmann in Inland Steel provided industrial
insights to the completion of my thesis. 1 must mention that Team Boyd aüows me ample
opportunities to interact among each other, and has been giving to me their cooperation
fiom time to time. Dr. Kevin Guangjun Cai helped me with thesis writing. Thanks are
atso due to the Ontario Centre for Materials Research for the financial support and Inland
and Stelco Steel companies for the materials supply.
Finaily, 1 certainiy tbank rny grandfather Hongxin Chen, my parents Xianming
Zhang and Rui Chen and my brother Gang Chen's family for their restless suppon that
encourages me to pursue my life goals. 1 will give my whole-hearted appreciation to both
Xiaofeng Zheng and Xiaolian Xu for ail these years' support: al1 afFections 1 could not
forsake.
This thesis is dedicated to Yuhua Luo, my grandmother who brought me up with
boundless love. Grandma in the sky, you are seeing me ciimbing to the height you wished
me to reach!
TABLE OF CONTENTS
.................................................................................. ABSTRACT i ACKNOWLEDGMENTS .............................................................. ii TABLE OF CON'IZNTS .............................................................. iii
........................................................................ LIST OF TABLES vii ... LIST OF FIGURES .................................................................... viii ... LIST OF SYMBOM .................................................................. XII[
ABBREVIATIONS .................................................................... xvi
CBAPTER 1 INTRODUCTION ....................................................................... 1
CHAPTER 2 EVOLUTION OF AUSTENI[TE CONDITION WITH TMP TIREATMENTS IN MICROALLOY ED STEELS ................................ 4 2.1 Chernistry of Austenite ................................................................. 4
2.1.1 Undissolved Precipitates ........................................................ 5 2.1 -2 S train-Induced Precipitates .................................................... - 7 2.1.3 Solutes ............................................................................. 9
2.1.3.1 Microalloyed Elernents in Solution ................................... 9 2.1.3 -2 Segregated MicroalIoyed Elements ................................. -10
2.2 Grain Structures of TMP-Processed Austenite .................................... -10 2.2.1 Recrystallized Austenite ....................................................... 1 1 2.2.2 Unrecrystailized Austenite .................................................... -13
CHAf TER 3 BAINITE TRANSFORMATION MECHANISMS Dl STEELS .............. 15 3.1 The Debate on Bainite Transformation Mechanism ............................... 15 3.2 Morphologies of Bainite .............................................................. 15 3 -3 Crystallography of Bainite Laths ................................................... -16 3 -4 Nucieation of Bainite ................................................................. -17
3.4.1 Bainite-Start Temperature .................................................... 17 3 .4.2 Nucleation Models ............ ................................................ 18
......................................... 3 .4.2.1 Classical Nucleation Theory 18 3 .4.2 -2 Olsen-Cohen Mode1 ................................................... 18
3.5 Growth of Bainite ..................................................................... 19 3 .5 . 1 Upper Bainite ................................................................... 21
. 3.5.1.1 Aaronson et al Theory ............................................... 21 3.5.1.2 B hadeshia-Christian Theory ......................................... -21
3 S.2 Lower Bainite ................................................................... 24 3 5 2 . 1 Spanos-Fang-Aaronson Mode1 ...................................... 24 3.5.2.2 Bhadeshia Theory ..................................................... 25
... I I I
CEiAPTER 4 EFFECTS OF AUSTENI[TE CONDITION ON BAINITE TRANSFORMATIONS IN MICRO ALLOYED STEELS ..................... 29 4.1 Definition of Bainite in Microailoyed Steels ....................................... 29
4.1.1 Low-C Steeis ................................................................... 29 4.1.2 Medium-C Steels ............................................................... 29
4.2 Recrystdized Austenite ............................................................. 30 4.2.1 Bainite Nucleation ................... .. ....................................... 30 4.2.2 Bainite Growth ................................................................. 30
4.3 Unrecrystallized Austenite ........................................................... 31 4.3.1 Bainite Nucleation ............................................................. 31 4.3.2 Bainite Growth ................................................................. 32
4.4 Cqstaliography ....................................................................... 32 4.5 The Effects of MicroaUoying EIements ............................................ 32
CBAPTER 5 ... EXPERIMENTAL ......................... .. .. .................................. 34
5.1 Materials ............................................................................... 34 5 -2 Therrnomechanical Processing Simulations ....................................... 34
5.2.1 Dilatometer Set-Up ............................................................ 35 5.2.2 Design of TMP Schedules .................................................... 36
5.2.2.1 Austenitization Temperature ......................................... 36 5.2.2.2 Austenite Conditions ................................................. -36 5.2.2.3 Coolhg Patterns ...................................................... -36
5.2.3 Determination of CCT Diagrams ............................................ -38 .................................................... 5.3 Microstructurai Characterization 39
5.3.1 Opticai and Scanning Electron Microscopy ................................ -39 .......................................... 5.3.2 Transmission Electron Microscopy -40
5.3.2.1 Specimen Preparation ................................................. 40 5 .3.2.2 TEM Examination .................................................... -42
...................................................................... 5.4 Hardness Testing 43
CEXAPTER 6 RESULTS ................................................................................. 44 6.1 Austenite Conditions ................................................................. 44
6.1.1 Evolution of Austenite Grain Structures .................................... -44 ........................................... 6.1.1.1 Austenite Grain Structures -44
6.1.2 Evolution of Precipitates Distribution in Austenite ........................ 45 6.1.2.1 L Steel ................................................................. -45 6.1.2.2 M Steel .................................................................. 51
6.2 Bainite Transformation Kinetics .................................................... -57 6.2.1 Microstructure-Based Definitions of Bainite ............................... -57 6.2.2 CCT Behaviour of L Steel .................................................... 58
6.2.2.1 L Steel As-Reheated ................................................ -64 ....................................... 6.2.2.2 L Steel Deformed Above Tm 65
6.2.2.3 L SteelDeformed Befow Tm ...................................... 65 ......................... 6.2.3 -4 L Steel Deformed Above and Below Tm 66
6.2.3 CCT Behaviour of M Steel ................................................... 66 6.2.3.1 M Steel As-Reheated ............................................... -66 6.2.3.2 M Steel Defonned Above Tm ...................................... 68 6.2.3.3 M Steel Deformed Below & ...................................... 71 6.2.3.4 M Steel Deformed Above and Below TM ......................... 72
6.3 Microstructures Observed by Scaaoing Electron Microscopy .................. -73 6.3.1 Austenite As-Reheated ......................................................... 73
6.3.1.1 L Steel .................................................................. 73 6.3.1.2 M Steel .................................................................. 77
6.3 -2 Austenite Deformed Above Tm .............................................. 79 6.3.2.1 L Steel .................................................................. 79 6.3.2.2 M Steel .................................................................. 82
6.3.3 Austenite Deformed Below Tm .............................................. 89 6.3.1 L Steel ..................................................................... 89 6.3.3.2 M Steel .................................................................. 90
............ 6.4 Microstructures Observed by Transmission Electron Microscopy 100 6.4. i Morphologies of Bainite ...................................................... 100
6.4.1.1 B: ...................................................................... IO6 6.4.1.2 B: ...................................................................... 109 6.4.1.3 BI ....................................................................... 110 6.4.1.4Bn ....................................................................... 111 6.4.1.5 Bm ...................................................................... 111 6.4.1.6 Brv ...................................................................... 111 6.4.1.7 a $ a w .................................................................. 112 6.4.1.8 a~ ....................................................................... 114
6.4.2 Intragranular Nucleation Sites of Bainite .................................. II5 6.4.2.1 Twin Boundaries ..................................................... 115 6 A2.2 Precipitates ........................................................... -119 6.4.2.3 Pre-Existing a* Laths ............................................... -119 6.4.2.4 Subgrain Boundaries ................................................. 123
6.4.3 Deformation-Induced Dislocations ......................................... -123
CHAPTER 7 ...... DISCUSSION ................... .........,...................... ................. 127
..................... 7.1 Evolution of Austenite Conditions with TMP Treatments 127 ..................................................... 7.1.1 Precipitates in Austenite -127
................................ 7.1.1.1 Types of Undissolved Precipitates 127 ............................. 7.1.1.2 Types of Strain-Induced Precipitates 128
..................... 7.1.1.3 Mechanisms of Strain-Induced Precipitates 130 ......................... 7.1.1.4 Kinetics of Strain-Induced Precipitation 131
7.1.2 Recrystallization of Austenite ............................................... 133
.................................................... 7.1.3 Substructure of Austenite -134 7.2 Variation in Bainite CCT Kinetics with Austenite Condition .................. 136
7.2.1 Bainite-Start Temperature .................................................. -136 7.2.2 Bainite Continuous Cooling Transformations ............................. 138 7.2.3 Deceleration of Continuous Cooiing Traw&onnation of Bainite ....... -139
7.2.3.1 Recrystallized L S tee1 .............................. .., ....... 139 7.2.3.2 Unrecrystaiüzed L Steel ............................................. 142 7.2.3.3 RecrystaUized M Steel .............................................. 142
7.2.4 Acceleration of Continuous Cooling Transformation of Bainite ........ 144 7.2.4.1 Uarecrystallized L Steel ................................. .. .......... 144 7.2.4.2 Unrecrystallized M Steel ............................................ 145
7.3 Bainite Nucleatim Sites ............................................................. 146 7.3.1 Grain Boundaries and Twin Boundaries .................................... L46
7.3.1.1 Recrystallized Austenite ....................................... .. ... -146 7.3.1 -2 Unrecrystaliized Austenite ......................................... -149
7.3 -2 Precipitates ..................................................................... 150 7.3.3 Pre-Existing Bainite Laths ................................................... 151 7.3 -4 Deformation Substmctures .................................................. -151
7.4 Growth of Bainite .................................................................... 153 7.4.1 Bainite Lengthening and Thickening ....................................... 153 7.4.2 Evolution of Bainite Morphologies ......................................... 156
7.4.2.1 Recrystallized Austenite ............................................ -159 7.4.2.2 Unrecrystaihized Austenite ......................................... -160
7.5 Bainite Transformation Models ................................................... -163 7.5.1 L Steel As-Reheated ......................................................... -163
7.5.1.1 Nucleation ............................................................ 163 7.5.1.2 Lengthening ............................................................ 163 7.5.1.3 Thickening ............................................................. 164 7.5.1 -4 Formation of Bainite Morphologies ............................... 164
7.5.2 L Steel Deformed Above Tm ................................................ 165 7.5.3 L Steel Deformed Below Tm ............................................... 168 7.5.4 M Steel As-Reheated ......................................................... 168
7.5 .4.1 Nucleation ............................................................ 170 7.5 .4.2 Lengthening .......................................................... -170
.... .. .................................... 7.5.5 M Steel Deformed Above TNR .. .. 170 7.5.6 M Steel Deformed Below T N ~ .............................................. 171
CEiAPTER 8 CONCLUSIONS ................... .. ................................................ 173
CHAPTER 9 SUGGESTIONS FOR FUTURE WORK ........................................ 176
CHAPTER 10 REFERlENCES .................... .................................. .............. 179
LIST OF TABLES
Table 2.1 Solubility Products for Some Cornmon Precipitates in Microaiioyed Steels. 5
Table 4.1 Variation of Bainite Crystaiiography with TMP.
Table 5.1 Chernical Compositions of L and M Steels.
Table 6.1 S v of Austenite Grain Structures.
Table 6.2 Number Density @) and Average Diameter (4 of Precipitates. 48
Table 6.3 Definitions of Continuously Transformed B ainite Types. 58
Table 6.4 Volume Fraction of Transformation Products in L steel. 64
Table 6.5 Volume Fraction of Transformation Products in M steel. 67
Table 7.1 Caiculated Solution Temperatures (Col,,) for Various Precipitates, OC. 127
Table 7.2 Calailated Precipitation Parameters of Nb(C,N) at 1 180 O C
for L and M Steels.
Table 7.3 Calculated 50% Recrystdization Time, b.3, of Austenite Deformed
Above TNR..
Table 7.4 Bainite Start Temperatures, OC.
Table 7.5 Cornparison of Interrupted TT Microstnictures of M Steel Defomed
Above and Below Tm
... in..
LIST OF FIGURES
Figure 3.4. 1 The Oison and Cohen mode1 for the development of a semicoherentbcc embryo fiom a perfèct screw dislocation [Olson and Cohen, 1976a-cl.
Figure 3.5.1 Free energy curves for a low (A), medium (B) and high (C) alloy steel showing the conditions necessary for the nucleation and growth of aw,
a~ and fE3hadeshia, 19921.
Figure 3.5.2 Schematic representation of Bm type upper a~ [Ohmori and M a h 199 11.
Figure 3-53 Sketch of the difFusionumtrolied rnechanism for lower Q formation [Spanos et aL, 19901.
Figure 3.5.4 Schematic representation of lower ae growth [Ohmori, l989].
Figure 5.2.1 Dilatometer set-up [Nelson, 19961.
Figure 5.2.2 Designed TMP schedules.
Figure 5.3.1 Schematic illustration of typicai dilatometer records on continuous cooling. (a-e) AL vs. T records for various cooling cycles; (f) cooling cycles for (a-e) nibscript "Y' is start temperature, "f' is finish temperature, T is temperature, L is length, t is time pldis, 19771.
Figure 5.3.2 Two methods of determining the transformation start and finish points.
Figure 6.1.1 TEM replica micrographs showing precipitate distributions in deformed+quenched L steel.
Figure 6.1.2 Precipitate size distributions in L steel.
Figure 6.1.3 (Nb,Ti)-rich particles nucleated on a MnS particle in L Steel defonned at 1000 O C .
Figure 6.1.4 TEM replica micrographs sliowing distributions of strain-induced precipitates in deforrned+quenched M steel.
Figure 6.1.5 Precipitate size distributions in M steel.
Figure 6.1.6 Three V-nch particles nucleated on a coarse @b,Ti)-rich particle in M steel deformed at 1025 OC.
Figure 6.2.1 ae morphologies-SEM: a) BI and Bma b) Bu and Bw, c) ~r~ and B:. 59
Figure 6.2.2 a~ morphologies-TEM: a) Br and Bma b) Bo and Bw, c) BP and BF. 60
Figure 6.2.3 CCT diagram of L steel. 62
Figure 6.2.4 AT-T curves at typical cooling rates of: a) 130, b) 10, and c) 1 OC/s. 63
Figure 6.2.5 CCT diagrams of M steel. 70
Figure 6.2.6 AT-T curves at typical cooling rates of: a) 5 OC/s and b) 1 "C/s. 7 1
Figure 6.3.1 SEM CCT rnicrostmctures - L steel as-reheated and cooled at: a) 130, b) 30, and c) 10 OCIs. 74
Figure 6.3.2 SEM IT microstmctures - L steel as-reheated and held at: a) 700 O C , 3600 seconds, b) 600 OC, 1800 seconds, c) 500 OC, 3600 seconds. 76
Figure 6.3.3 SEM CCT micros~nictures - M steel as-reheated and cooled at a) 10, b) 5 , and c) 1 O C k . 78
Figure 6.3.4 SEM CCT microstructures - M steeI as-reheated. a) a0 nucleated at Nb-nch precipitate, b) The corresponding EDS spectrum for Nb-rich precipitate. 79
Figure 6.3.5 SEM IT microstructures - M steel as-reheated and held for 1800 seconds at: a) 600, b) 500, and c) 400 O C . 80
Figure 6.3.6 SEM CCT microstructures - L steel defomed at 1000 OC and cooled at: a) 30, b) 10, c) 5, and d) 1 "C/s. 83
Figure 6.3.7 SEM interrupted CCT microstmctures - L steel deformed at 1000 OC and: a) 10 *C/s to 500 O C , quench, b) 10 OC/s to 550 OC, 1 OC/s cool, c) 5 OC/s to 520 OC, quench. 84
Figure 6.3.8 SEM IT microstmcture - L steel deformed at 1000 OC and held for 1800 seconds at 400 O C . 84
Figure 6.3.9 SEM CCT microstmctures - M steel defomed at 1080 OC and cooled at: a) 10, b) 5 , and c) 1 "C/s. 85
Figure 6.3.10 SEM CCT microstructures - M steel deformed at 1025 OC and cooled at a) 10 and b) 5 "C/s. 86
Figure 6.3.1 1 SEM intempted CCT microstructure - M steel deformed at 1 080 OC and cooled at 5 OC/s to 520 OC, quench. 86
Figure 6.3.12 SEM IT microstructures - M steel deformed at 1080 "C and held at: a) 600 OC, 1800 seconds, b) 500 OC, 5 seconds, d) 400 OC, 1800 seconds, e) 400 OC, 5 second. 88
Figure 6.3.13 SEM CCT microstructures - L steel deformed at 850 OC and cooled at: a) 30, b) 10, c) 5, and c) 1 "C/s. 92
Figure 6.3.14 SEM CCT microstructures - L steel defomed at 780 OC and cooled at: a) 30, b) 10, c) 5, and c) I "C/s. 93
Figure 6.3.15 SEM CCT microstructures - L steel deformed at 1 O00 and 780 OC and cooled at: a) 30, b) 5, and d) 1 "C/s. 94
Figure 6.3.16 SEM IT microstmchires - L steel deformed at 780 OC and held for 1800 seconds at: a) 600, b) 500 OC. 95
Figurl: 6.3.17 SEM CCT microstructures - M steel deformed at 950 O C and cooled at: a) 10 and b) 5 "C/s. 96
Figure 6.3.18 SEM CCT microstnichires - M steel deformed at 850 O C and cooled at: a) 10 and b) 5 "C/s. 97
Figure 6.3.19 SEM CCT microstructures - M steel deformed at 1080 and 850 "C and cooled at: a) 10 and b) 1 *C/s. 98
Figure 6.3.20 SEM intempted IT microstmcnires - M steel as-reheated and held for 5 seconds: a) 600, b) 500, c) 400 OC. 99
Figure 6.4.1 L steel deformed at 1000 O C and cooled at 10 OC/s-Overall morphology. 10 1
Figure 6.4.2 L steel deformed at 850 O C and cpoled at 10 OC/s-OveralI morphology. 102
Figure 6.4.3 L steel deformed at 780 OC and cooled at 30 *Us-Overall morphology. 103
Figure 6.4.4 L steel deformed at 780 OC and cooled at 10 *Ch-Overall morphology. 104
Figure 6.4.5 M steel deformed at 1025 OC and cooled at 5 OC/s-Overall morphology. 105
Figure 6.4.6 M steel deformed at 850 OC and cooled at 3 OC/s-Overall morphology. 105
Figure 6.4.7 L steel as-reheated-B:.
Figure 6.4.8 L steel as-reheated-B:.
Figure 6.4.9 L steel as-reheated-Br.
Figure 6.4.10 M steel as-reheated-Bu.
Figure 6.4.1 1 L steel as-reheated-Ba.
Figure 6.4.12 M steel as-reheated-Bw.
Figure 6.4.13 M steel as-reheated-GBa-nucleated a w / a B .
Figure 6.4.14 M steel deformed at 1025 OC and cooled at 5 OC/s-Twin boundary- nucleated aw/aB.
Figure 6.4.15 m. a) L steel as reheated - 30 "Ch. Auto-ternpered au laths with multi-variants of 0.
b) M steel as-reheated - 30 OCfs. a~ plates.
Figure 6.4.16 Nucleation of a~ at twin boundary.
Figure 6.4.17 Nucleation of aB at precipitate.
Figure 6.4.18 Nucleation of as at pre-existing ae plates- M steel deformed at 850 OC and cooled at 5 "C/s.
Figure 6.4.19 Nucleation of aB at subgrain boundary- L steel deformed at 850 OC and cooled at 10 "Cfs.
Figure 6.4.20 Bm with fuzzy lath boundaries and ellipsoidal intraiath 8- L steel deformed at 1000 OC and cooled at 10 W s .
Figure 6.4.2 1 Tangled disiocations in ~ 1 ~ -
L steel defonned at 780 "C and.coo1ed at 30 " C k
Figure 6.4.22 Low density o f dislocations in BF- L steel deformed at 780 OC and cooled at 1 OC/s.
Figure 7.1.1 5% precipitation tirne ( t o . ~ ) vs. defomation temperature ( T d ) .
Figure 7.3.1 GBa-nucleated a~ - M steel deformed at 1025 OC and cooled at: a) 10 and b) 5"C/s.
xii
Figure 7.4.1 The variation of bainite dimensions with cooling rate in L steel as-reheated.
Figure 7.5.1 Nucleation of bainite in L steel as-reheated at: a) intermediate cooling rate, b) high cooling rate, c) slow cooihg rate. 1 66
Figure 7.5.2 Lengthening of bainite in L steel as-reheated. 166
Figure 7.5.3 Thickening of bainite laths. 166
Figure 7.5.4 Formation of bainite morphologies. 167
Figure 7.5.5 Nucleation of bainite in L steel deformed below Tm. 169
Figure 7.5.6 Nucleation of bainite in M steel as-reheated. 171
Figure 7.5.7 Lengthening of GBa-nucleated bainite in M steel as-reheated. 172
Figure 7.5.8 Formation of BN in M steel as-reheated. 172
LIST OF SYMBOLS
constant. constant. carbide-fiee bainitic femte laths with retained austenitdmartensite films. carbide-fiee elongated bainitic femte laths with retained austenitdmartensite films. granular bainitic femte laths with retained austenitdmartensite islands. bainitic ferrite laths with interlath carbides. bainitic ferrite laths with intraiath carbides. bainitic ferrite laths with both intra- and interlath carbides or austenite/martensite. bainite finish tempera- Iower bainite bainite start temperature. upper bainite constant. constant. solute concentration in eqwlibrium with a particle of idmite radius. precipitate mean diameter. weighted average difisivity of carbon in austenite. carbon diffisivity in austenite. initiai austenite grain diameter. recrystallized austenite grain diameter. solute diasion coefficient. equilibrium subgrain diameter. austenite grain diameter. constant. precipitate volume fraction. constant. constant. constant. ferrite start temperature. bainite growth rate. activation energy for bainite transformation. grain boundary ferrite. bainite growth rate in carbon steels. universai nucleation finction. fùnction representiog the critical value of the fie energy change needed before the athermal, diffusiodess nucleation and growth of becornes possible. strain energy. stored energy of (- 400 J/mol). stored energy of a w (- 50 Jhol). chernical driving force for nucleation of Nb(C,N). driving force for bainite transformation. released fiee energy due to the destruction of a defect. Free energy change for heterogeneous nucleation.
total free energy for precipitation. misfit strain energy. volume fkee energy. grain size heterogeneity factor. bainite nucleation rate per unit volume of austenite. coefncient representing the effects of m] on the suppression of bainite growth. coefficient representing the effects of m] on the suppression of bainite growth. saturation ratio, achiai amount in solution to equilibrium amount in solution. static coarsening rate constant. dilatometric sample length change. original dilatometric sample length. deforrned dilatometric sample length. ratio of half-width to length of bainite lath. final dimension of bainite. concentration of microaUoying element Mi in austenite. equiiibrium concentration of microdoying element Mi in austenite. martemite start temperature. martemite finish temperature. p earlit e . pearlite s t a t t emperature . austenite grain radius. cntical particle size above which grain growth occurs. particle radius at time O. particle radius at t h e t. gas constant. surface area. time. 5% precipitation tirne. 50% recrystallization tirne. isothermal holding time. bainite incubation tirne. time to decarburise the carbon-supersaturated femte. temperature. deformation temperature. top temperature for bainite C curve on isothermal transformation diagram no-recrystallization temperature. . temperature where Gibbs fiee energy of the unstressed austenite is equai to that of femte of the same composition. precipitate solution temperature. volume of femte. precipitate molar volume fiaction. average mole fi-action of carbon in the aiioy. paraequilibrium carbon concentration in femte. paraequilibnum carbon concentration in austenite. bainite volume fiaction.
wdmanstatten ferrite start temperature. Zener-Holloman parameter. femte. bainite. lathlike bainite. lathless bainite. ruartensite. W~dmanstatten femte. austenite. particle/matrix interfacd energy. retained austenite. tme strain. cnticd straio for dynamic recovery.
strain rate. iron carbide. transational ïron carbide. precipitate number density. interfacial fiee energy. width of bainite lath.
bcc body centered cubic. CCR conventional controiied rolline.
Y
CCT continuous woling tra&ormation COR crystaiiographic orientation relationship. EDS energy dispersive spectroscoov. * d
fcc face centered cubic; GB grain boundary. LVDT iinear voltage differential transducer. N-W Ni~shiyama-Wassemian orientation relationship. IT isothermd traasfonnation. K-S Kurdjumov-Sachs orientation relationship. RCR recrystallized contr011ed roiiing. SADP selected area diEraction pattern SEM scanning electron microscopy. SDLE solute drag-like effect. TEM transmission electron microscopy. TMP thermomechanical processhg.
INTRODUCTION
Developed over 35 years ago Borchynsky, 19771, microalloyed steels cm be
defined as C-Mn steels containing s d amounts (usudy l e s than 0.15 W.%, singiy, or
in combination) of Nb, V or Ti (Pauies, 199 11.
By microdoying, the microstructure can be rehed and the strength can be
enhanced by the precipitation of smaü microaiioy carbides (NbC, VC, Tic), nitrides (MN,
VN? TN), or complex carbonitrides (Nb(C,N), V(C,N), Ti(C,N)) paules, 19911. Further,
a tailored thermomechanical processing (TMP) schedde and controlled cooling can obtain
strength Ievels equivalent to those in quenched+ternpered steels, but at a substantial
economic advantage as the heat treatment cycle is eiiminated [Pickering, 19831.
TMP-processed low-C bainitic steels have higher strength and toughness than
those of conventional feriite+pearlite (a+P) steels [Garcia et al., 19901. Recently ,
microalloying (usuaiiy - 0.1% microalloying elements such as V, Nb and Ti [Kuziak and
Cheng, 199 11) was successfidly applied to medium- to higher-C (up to 0.8%) bar and rod
steels thermomechanicaily processed at higher temperatures than strip and plate steels
paules, 199 1; Krauss, 19891. For forging steels, an a+P structure is desired in the low
strength range (350-450 MPa) Pickering, 1983; Jonas et al., 19851, while bainite (a*) is
usuaiiy considered detrimental to both mechanicd properties and machinability
warchwica et al., 19931. In the higher strength steels (> 600 MPa), fully bainitic or
predominately baintic microstructures are desued. As austenite (y) condition (y grab
structure and substructue, solute and precipitates in y) determines ae transformations, the
critical problern is to develop TMP scheduies for desired y condition in order to control
the development of as structures. However, there are few systematic studies on the
effects of y conditions on as transformations in microdoyed steeis, especidy during
continuous cooling trdormation (CCT) that is always associated with industrial
practice.
'Bainite" was named in 1934 in honour of E. C. Bain who, together with E. S.
Davenport in 1925, had found an "aciailar, dark etching aggregate'"ch f o m &ring y
is0tfienr.d decomposition at temperatures above that at which martede (m) nrst for- but
below that at which fine P is found pavenport and Bain, 19301. Two schools ofthought have
existed regardhg whether the ors transformation mechanism is diffusion-controiied or
displacive in nature wtemational Cod. on Bainite, 19901. The &sional opinion
considers that ae lathdplates develop by a diffusional ledge mechanism, while the
displacive opinion claims that as IaWpIates grow without diaision of solvent and
substitutional solute atoms.
There is even more confusion in interpreting ae microstnictures in modem
microalloyed steels controlled cooled after hot roiling because UB microstnictures are
usually too complex to be denned in the conventionai way m o n d s and Cochrane, 1990;
Araki and Enomoto, 19901.
Therefore the objectives of the present research are:
1) to define the complex a~ microstructures in microalloyed bar steels containing
Iow- and medium-C, respdvely, and cooled continuously,
2) to investigate the effects of y condition on the ae transformation kinetics and %
morphologies, and
3) to develop an a~ transformation mode1 reflecting the influences of various y
conditions.
It is believed that this research can assist the clarification of the complex a~
microstructures, provide insights to nucleation and growth mechaaisms of a~ and the
effects of y conditions on a~ transformation kinetics and morphologies, and thus benefit
steel companies in controlliug as microstructures during industriai procusing.
EVOLUTTON OF AUSTENXTE CONDITION WïMt TMP TREATMENT
IN MICROALLOYED STEELS
TMP is a technique designed to improve the mechanid properties of materials by
controiiïng the hot deformation processes, which were originally designed to produce the
extemal shape of products [Sekine, 1988; Jones et al, 19851. Using microdoying and
designed TMP scheduies, the y condition of microaiioyed steefs can be tailored to obtain
designed microstructures.
2.1 CEWMISTIRY OF AUSTENITE
It has been weii established that microailoying causes a remarkable retardation of
the restoration foiIowing deformation due to 1) solute-drag Weiss and Jonas, 1979;
Lamberigts and Greday, 19771, 2) strain-induced precipitation of fine Nb(C,N) [Irani et
al., 1967, Sekine and Maruyama, 1976; hine and Baker, 19791, or 3) the combined
eEects of solute Nb and precipitation b b e r i g t s and Greday, 19771. The main effects of
microalloying elernents such as Nb and Ti are to suppress the y grain coarsening and raise
the no-recrystdization temperature (Tm ) [Palmiere et al., 19921. Therefore, the
chemistry of y (precipitates and solutes) determines the restoration behaviour of deformed
Y-
2.1.1 Undissolved Precipita tes
At equilibrium, the arnount and composition of precipitates are determined by their
corresponding solubility at the reheat temperature. Solubility products of some cornmon
precipitates in microalloyed steels are summarized in Table 2.1.
Table 2.1 Solubility Products for Some Cornmon Precipitates in Microailoyed Steels
Solubility Products Reference Irvine, 1967 Narita, 1975 Narita, 1975
Houghton et al., 1982 Houghton et al., 1 982
Narita, 1975 f i n e , 1967
Fountain and Chipman, 1 992
At high reheat temperatures, there are few undissolved precipitates in y, and the
suppression of the growth of y grains is mainly attributed to solute drag effects of Nb. Ti
and V. However, at low reheat temperatures, the undissolved fine precipitates with low
solubility products such as TiN, Ti(C,N) and Nb(C,N) or their compounds apply an
effective constraint to pin the y grains. Hence the average y grain diameter D, can be
estimated by Zener's relation [Zener, 19491
where r is the average precipitate radius. and f is the volume fraction of precipitates.
There is a cnticd pinning particle size, r,, above which the grain boundaries
become unpimed and grain growth occurs [Gladman, 19661
where H is the grain size heterogeneity factor.
During reheat, precipitates in y usually coarsen following the Lifshitz-Wagner
equation [Lifshitz and Slyozov, 196 1 ; Wagner, 196 1 ]
and
where r, and r, are the particle radii at time t and 0, respectively, K is the static coarsening
rate constant, Ce is the concentration of solute in equilibrium with a panicle of infinite
radius, y, is the particiehatrix interfacial energy, V, is the molar volume of precipitate, D,
is the solute diffision coefficient, and c is a constant.
If the coarsening mechanism is interface reaction controlled, n = 2; buik difision
controlled, n = 3 [Wagner, 19611; grain boundary diffision controlled, n = 4; and
dislocation pipe diffusion controlled, n = 5 [Ardell, 19721.
2.1.2 Strain-Induced Precipitates
lo microaiioyed low-C steels, Nb(C,N) is the most common strain-induced
precipitate. Complex precipitates of cubic structure such as (Nb,Ti)-, (Nb,V)- and (Ti,V)-
nch particies are also found [Loberg, et al., 1984; Crooks et al., 198 1 ; Suzuki et al.,
19871. In microalloyed medium-C steels, the common precipitates in y are Nb(C,N) or
cornpounds of alloy elements such as (Nb,Ti)- and (Nb,V)-rich particles that control the y
grain size during reheating, and V(C,N) that forms in a for precipitation strengthening
purposes [Prikryi et al., 1994; Paules, 19911. For simplicity, only the effects of the
precipitation of Nb(C,N) is considered.
Nucleation and growth of strain-induced Nb(C,N) precipitates are considered to
occur at: 1) y grain boundaries [Palrniere et al., 19921, 2) dislocation nodes in the three-
dimensional network of dislocations and 3) subgrain boundaries generated by deformation
[Dutta et al., 19921.
The free enthalpy change for precipitation to occur is the saturation ratio k,
defined as the ratio of the actual amount in solution to the equilibnum amount at
temperature T .
where p l , [Cl and M are the concentration of Nb, C and N in y.
Then the chernical driving force for nucleation of Nb(CN) is:
where [NbT, [CY and are the equilibrium concentrations.
The total fiee energy change for precipitation is:
Based on nucleation theory and empincal anaiysis, an equation to estimate the time
for 5% precipitation of Nb(C,N), h m , has been proposed pu t t a and Sellars, 19871
t , , , = A [ N ~ ] - ' E - 'Z -" ' 270000 B
RT T 3 (ln k, ) ' (2.8)
where E is the true strain of the prior deformation carried out at a strain rate E and
deformation temperature T d leading to the Zener HolIomon parameter,
. Z = E- exp
400000 , A and B are constants.
R T d
The time for the completion of precipitation is about two orders of magnitude
greater than io.05 [Dutta, Vaides and Seiiars, 1992 1.
2.1.3 Solutes
2.1.3.1 Microalloying EIements in Solution
For hot deformation, the no-recrystaiiization temperature (&) is the temperature
below which complete recrystallization of y grains does not occur between deformation
steps.
Nb. Al, Ti, and V have been shown to retard y recqstallization [Irvine et al., 1970;
Cuddy, 198 11. For instance. the retardation of recrystaliization due to Nb occurs only
when Nb is in solution in y before deformation. If Nb rernains undissolved because of a
low reheat temperature, it does not demonstrate any delaying effect [Irvine et al., 19701.
For Nb-bearïng steels, soiute Nb atoms retard recovery and recrystallization until the
occurrence of strain-induced precipitation, while strain-induced precipitates retard the
onset and progress of recrystallization [Tanaka, 19881. In ternis of this concept, a
regression equation has been developed by Boratto et al. [1988] to estirnate T M R .
The equation is valid in content ranges (wt-%): 0.04 < C 0.17, 0.15 < Si c 0.50,
0.002<Al<0.65,0.41 Mn< 1.9,O<Nb<0.06,O<V<0.12, O<Ti<0.11, O<Cr<0.67,
O < Ni < 0.45.
2.1 -3 -2 Seareeated Microdoyina Elements
During reheat, it is well known that some elements such as B and Nb retard y
recrystallization and suppress the formation of both a and as by segregating to y grain
boundaries thus reducing the interfacial energy welloy et al., 1973; Jung et al., 19951.
There are two kinds of non-equilibrium segregation in steels containing Nb+B.
The first takes place at y grain boundaries and is attributable to the excess vacancies
produced by the deformation, which delays the recrystallization start time to a limited
degree. The second takes place possibly by forming (Nb,B)-complexes at the "fiesh
grain boundaries produced by recrystallization [He et al., 19881. In die case of steels
containing Nb+Ti, Nb and Ti are able to either segregate to y grain boundaries, or
precipitate as extremely fine carbides distributed aimost continuously at y grain boundaries
with a specific orientation relationship with y [Sharma and Purdy, 19731.
2.2 GRAIN STRUCTURES OF TMP-PROCESSED AUSTENITE
During and afler hot deformation, four possible restoration processes occumng in
work-hardened y are 1) dynamic recovery and 2) dynamic recrystallization thai occur
during deformation, and 3) static recovery and 4) static recrystallization that occur fier
deformation. resulting in soflening of work-hardened y Parr and Tipper, 1947; Sekine,
19881. TMP above TNR produces recrystallized y in which dynamic or static
recrystallization is possible, while TMP below Tm produces unrecrystallized y in which
only recovery could occur.
2.2.1 Recrystallized Austenite
Dynamic recrystaiiization occurs when the applied sixain ( E ) exceeds a critical
value. When the accumulated strain is smdl, only dynamic recovery may take place
[Tamura et al., 19881. For Nb-bearing steels, the critical strain (EJ for dynamic
recrystallization is related to the initial y grain diarneter (Do, pm) and the Zener-Holloman
parameter (2, 1 /s) [Sellars, 1980; Dutta and Sellars, 19871
where E = 6.4 x 104 for Nb-bearing steels.
The smallest E, is obtained when the sarnple has the srnailest D,, and is deformed at
s
highest Td and the lowest E . However, for mmt commercial steels, it is difficult to obtain
dynamic recrystallization during normal roliing schedules as E, is very large WcQueen,
19681. Even dynamic recovery is only observed in steels deformed in the a region [Ouchi,
When strain exceeds the cnticai value for static recrystallization to occur
(determined by the prior deformation conditions and the pnor grain size), static
recrystallization starts in work-hardened y by nucleation of new grains predominantly at
triple junctions of y grains and grain boundaries rather than grain intenors wozasu et a/.,
197 1 1. The progress of recrystallization is essentially the migration of the recrystallizing
front into the deformed matnx. Increasing the holding time at temperatures where static
recrystallization occurs will develop annealing twins gradually [Tamura, 1 9881. The
completion of recrystalhtion is followed by normal grain growth [Tanaka, 198 11.
The staticaily recqstallized y grain diameter, D r , is related to Do and E by the
following equation [Sellars, l98O]
For Nb-treated steels, the time of 50% recrystallization (kJ) is given by [Sellars,
19803
Io, = 2.52 x 1 0 ' ' ' ~ ~ ~ ~ exp 325000
RT
130000 IO s = 9.24 x I O - ~ D , E exp
RT
In industrial rolling, recrystallized y is obtained in Recrystaiiizaîion-Controlled
Rolling (RCR) that involves control of y grain growth during reheating, repeated
deformation above Tm, and inhibition of y grain growth dunng and afler rolling [Paules,
199 11. Dunng forging of Nb-bearing steels, hi& reheat temperatures and high T d oAen
result in a large D,, so Equation 2.8 can be used to predict D, with good accuracy
wuziak and Cheng, 1992).
2.2.2 Unrecrystallized Austenite
There are several microstructurai changes in unrecrystallized y due to deformation
below Tm. First the ratio of the surface areas before : after rolling is largely increased
due to the change of undeformed spherical y grains to pancaked ellipsoidal y grains, e.g.,
the y surfaces area increases by 25% for 50% deformation [Tamura, 19881. Second,
ledges (or steps) on the y grain boundaries and a high dislocation density near deformed y
grain boundaries are introduced. Third, incoherent annealing twins with a large number of
ledges on the twin boundaries are produced [Tamura et ni., 1988; Amin and Pickering,
198 11. Last, deformation bands are generated, which appear as closely spaced parallel
lines and often teminate within a grain [Kozasu et al., 19771. The regions of deformation
bands are characterized by a high density of cells consisting of tangled dislocations.
Usually the deformation band density is little affected by Td in the non-recrystallization
region. With the increase of reduction, deformation band density increases rapidly
[Kozasu el al., 19771. Therefore, the free energy of unrecrystallized y is higher than that
of recrystallized y due to the stored energy, which is a Function of dislocation density.
Dislocations are potential nucleation sites for new phases through its stress field [Tamura,
19881.
After deformation below Tm, s m d subgrains with dense subboundaries usually
fonn in y-Fe due to its rather low stacking fault energy [McQueen and Jonas, 19841. It
was reported that a dispersion of fine precipitates (e-g., strain-induced precipitates) c m
nucleate at subboundaries, thus stabiiizing the substructure and determinhg its scale
[Akben et ai., 198 1 ; Jonas and Akben, 198 1 ; Oblak and Owezarski, 19721. Both soiutes
and particle dispersions inhibit static recrystdlization thus making it easier to preserve the
hot work substmctures dunng cooling WcQueen and Jonas, 19751. The equilibrium
subgrain diameter, LIN6, can be represented as
where FI, F;, and F3,are empirical constants [McQueen and Jonas. 19841.
In industrial processing, the conventional controlled rolling (CCR) or controlled
forging involves rolling or forging at temperatures below TM to obtain refined y grains
[Tamura, 19881.
BAINITE TRANSFORMATION MECBANISMS IN STEELS
3.1 TBU DEBATE ON BAIMTi3 TIRANSFORMATION MECaANTSM
Two opposing theories have been proposed for the growth of bainitic a durlig a~
transformations. The displacive theory considers that the atornic rearrangements during
bauiitic a growth occur by a difiùsionless shear mechanism as far as the substitutionai atoms
are concernai, although the diffùsion of interstitiai atoms such as C is aiiowed Phadeshia and
Chnstiaq 19901. The diaisional theory considers that the f i s i o n of substitutional atoms
during bainitic cr growth is essential in the vicinity of the advancing aly interfaces [Hehemarüi
ef al., 1 972; Rigsbee and Aaronson, 1 9791.
3.2 MORPHOLOGIES OF BAINITE
The description of the two major a~ morphologies, upper and lower as, was first
used by Mehl [1939] to distinguish between morphologies of ae formed at higher and
lower temperatures, respectively. Upper a~ consists of an array of a Iaths or plates and
interlath cementite (8) layers [Shimizu et al., 19641. Lower a* is composed of sheaves
(bundles) of a plates in which 0 platelets are oriented at an angle of approximately 55 to
60 O to the plate axis mumg and Thomas, 1977; Lai, 19751,
Ohmon et al. [1971] tenned the different types of isothermally transfomed upper
a~ as BI: 0-free ae laths; Bu: a* Iaths + interlath 8; and Bm: a~ taths + intralath 8. Others
classified Bm as lower ae pramifitt and Speer, 19901, while some insisted that Bm should
be upper ae due to its lathlike morphology rather than the platelike lower a~ [Ohrnori and
Maki, 1991 1. A structure featuring both inter- and intralath 8 was defined as lower ae
phadeshia, 1980; Ohtani et al., 1 WO]. In the current study, the nomenclature proposed
by Ohmon has been used for descnbing the various as morphologies (see Section 6.2).
3.3 CRYSTALLOGRAPEW OF BAINITE MTHS
For a bainitic a lath idealised as a parallelepiped with dimension a > b z c, the
crystaiiography has been characterised in detail by Davenpon [1974] as follows: Growth
direction: [i0 11, [([;II, ; Habit plane (ara = ab): (232) @4) : Face of a r a (ac): ( 1 0 1 ) ~ ; r
Orientation relationship: K-S (Kurdjumov-Sachs): [i0 l ] ,~ ( [%J, and ( 1 1 l)? 11 (O 1 Ila .
Hence, the major growth direction of each lath is the close-packed direction of the bainitic a
and y lattices. Sandvik [ l98ZaI reported an orientation relationship of (1 1 1 X (1 (0 1 1), and
[IO 11, approximately 4' from [fil], , which is close to the N-W (Nishiyama-Wasserman)
relationship.
Generaily, the habit plane of lower ae is considered irrational. In an Fe-Cr-C
alloy. the habit plane of lower a~ lies close to (254h and the orientation relationship
between bainitic a and y is near to the K-S relationship [Srinivasan and Wayman, 19681.
3.4.1 Bainitestart Temperature
The a-start temperature (B3 can be defined as the highest temperature at which ae
phase is observed to form at a detectable rate phadeshia, 19921.
Foiiowing Bhadeshia et al. [Bhadeshia and Edmonds, 1980; Bhadeshia and Waugh.,
198 11, a* occurs below Ta a temperature where the Gibbs fke energy of the unstressed y is
quai to that of a of exady the same compositioa This is supporteci by the observation that
the C concentration in untransformeci y is close to the T, iine in the incomplete c c ~
transformation.
On the reconstructive side, the incomplete as transformation is caused by a strong
solute drag-like effect (SDLE). Below B,, as nucleates at y grain boundaries. The SDLE
Lirnits the extent to which the bainitic a grows, and new a crystals (sympathetically) nucleate at
the immobW a /y boundaries. This renucleation process conrinues until the çurrounding y
becornes suflïciently e ~ c h e d Ui C to prevent hrther nucleation of bainitic CI, leading to
transformation stasis [Aaromon et al., 1 9901.
Steven and Haynes [1956] developed an empUical equation for B, that is valid in the
following composition range (wt%): 0.1 -0.55C%, 0.0-5 .ON%, 0.1-0.3 5Si%, 0.0-3. SCPh, 0.2-
1 .7Mn%, and 0.0- 1 .OMO.
Bs ("C) = 830 - 270C - 90Mn - 37Ni - 70Cr - 83Mo
3.4.2 Nucieation Modeis
3.4.2.1 Classical Nucleation Theory
Classical nucleation theory is based on the occurrence of random phase fluctuations in
the parent phase [Christian, 19751. This model must be regardai as a reconstructive nucleation
process. The activation energy, 6, has an inverse square relationship to the driving force,
AGch, such that : G' = AG^'*.
3 -4.2.2 Oison-Cohen Mode1
As shown in Figure 3.4.1, the model proposed by Olson and Cohen [ 1976a-c] assumes
that the embryonic defects are closely-spaced group of faults derived fiom the dissociation of
other defects already present in y. Figure 3 -4.1 a shows fcc y, and Figure 3.4.1 b shows the
three dimensionai dissociation of a dislocation over a set of three close packed planes. The
structure thus produced is not yet bcc. Figure 3 . 4 . 1 ~ shows the relaxation of the fault to a bcc
structure, involving the introduction of partial dislocations in the interface. Figure 3.4. Id
shows the addition of perfkct screw dislocations which cancel the long-range main field of the
panid dislocations introduced in Figure 3 -4. lc.
As the temperature and fault energies fa the embiyo develops atherrnaiiy into a thin
plate of a ~ , which might subsequently thicken by some self-reproducing rnechanism such as
the pole mechanism proposed by Christian [1975]. Thermaüy activated nucleation then
corresponds to the attempts by the emb~o/matrÎx interface to overcome an interface fiction
stress. G' bas b e n found to be dûectly proportional to AGch rather than the inverse square
relationship as show in the dassicai nucleation theory G' Gmm + AGc.. Here G,, is the
strain energy.
a bcc -[O 1-11 8
afcc - a bcc - -[1 1 O] = -[1 1 1) 2 2
Figure 3.4.1 The Olson and Cohen mode1 for the development of a semicoherent bcc embryo fkom a perfiect screw dislocation [Olson and Cohen, 1976a-cl.
Bhadeshia and Christian [1990] considered that there is a cornpetition for the growth of
a w (Widmanstatten a), a~ and m. For example, the nucleation process is identical for a~ and
a ~ . if at the top temperature, Th, of the as "Cu curve on the IT ( i so thed transformation)
diagram, the dnving force avaiiable is nitncient to account for both diffusionless transformation
and the stored energy of ae, a~ does not fom. Otherwise, crw foms and Th = W, (a* %art
temperature).
The nuclei that can M y evolve into a new phase can be schematically show in
Figure 3 S. 1 of the fiee energy w e s of three different aeels containing increasing quantities
of y-stabiiising elements. Gw is the stored energy of a w (- 50 J/mol), Gm is the stored energy
of a~ (- 400 J/mol) and G ~ O ' is the bction representing the critical value of the f i e energy
change needed before the athennai, difisiodess nucleation and growth of becomes
possible, which is relatively insensitive to solute concentration [Bhadeshia, 198 1bJ. GN is the
so-caiied universal nucleation fùnction applicable to ali [ow aüoy steels [Aii and Bhadeshia,
1 9901 :
In A (low alloy steel), ail three transformations (aw, aa and ahr) are expected as the lT
temperature is reduced. For B (medium alloy steel), at the temperatures where a~ nucleation
becomes possible, the growth condition for ae is aiso satisfied, so that any nuclei evolve into
a ~ . In C (hi& alloy steel), a w and are eliminated, and only -1 forms.
3.5.1 . 1 Aaronson el al. Theow
From a cifisional mechanism standpoint, ae grows by forming growth iedges in
paniaiiy coherent a/y boundiuies to overcome interfacial structure barriers [Aaronson, 1969;
Aaronson et d, 19701
The upper a~ laths form by a face-to-face sympathetic nucleation [Aaronson and
Wells, 19561, containing 8 predorninately nucleated at the broad faces of the individual a laths
and growing preferentidy dong hem, thus causing 0 to be pardei to the long axis of the laths
[Oblak and Hehemann, 19671. Sympathetic nucleation is the nucleation of a precipitate crystal
at the interphase boundary of a previously formed crystal of the same phase when the rnatrix
and the precipitate dzer in composition [Aaronson and Wells, 19561. a~ plates lengthen and
ducken at diffusion-controlled rates. The a/y boundaries associated with proeutectoid a and
upper a* plates are sessile and hence unable to move conservatively by glide [Rigsbee and
Aaronson, 1979; Li elal., 19881.
3.5.1 -2 Bhadeshia-Christian Theory
In the displacive mechanism view, laths ofaB nomally nucleate at y grain boundaries
and propagate toward the grain interiors by the nucleation and growth of individual subunits
with new subunits nucleating near the tip of a previous subunit. Both upper and lower aa
consia of aggregates of laths or plates of bainitic a separated by regions of residual phases
phadeska and Christian, 19901. The bainitic a focms without any diffusion initiaily and then
C atoms are rejected from the C-supersaturated bainitic a into the parent y by diffusion,
resuiting in the e~chrnent of C atoms in the parent phase [Bhadeshia, 198 1; Bhadeshia and
Waugh, 19811.
TEMPERATURE -+
Figure 3.5.1 Free energy curves for a low (A), medium (B) and high (C) alloy steel showing the conditions necessaq for the nucleation and growth of UN-, c t ~ and cy\l
hadeshia, 1 992 1.
Bhadeshia [1985] demonstrated that in plain C steeis the rneasured lath growth rates
are much higher than would be indicated by C difision-controlled growth. Also, the growth
rate of individual subunits of ae has been measured to be orders of magnitude faster than
would be expected from C difision-controlled growth for a Fe-Mn-Si-C alloy [Bhadeshia,
1 9841.
Ohtani et al. [1990] proposed a mode1 to describe the displacive formation of Bm Li
upper a~ (Figure 3.5 -2). The bainitic a needla having the pardeIogram cross sections sirnila
to BI nucleate initidy, and reject the supersaturateci C atoms into the y (Figure 3.5.2a). C
atoms enriched locdy in y in contact with the a needle sides, which correspond to the lamce
invariant shear planes, precipitate as 6 platelets at these intefiaces probably due to the fiict that
the d y lattice matching on these interfàces is not so coherent as that on the habit planes of the
other sides and induces higher intefice energy (Figure 352b). The side-by-side formation of
these a~ subwiits on { 1 1 1 }, - {575), lads to the lathlike a~ involving the û platelets aligning
on a specific crystaliographic plane (Figure 3 .52~) .
Figure 3.5.2 Schematic representation of BWtype upper a~ [Ohmon and Maki, 199 11. a) a subunit formation, b) the precipitation of 0 platelets on the side suifaces of the subunit, c) the codescence of the subunit into Bnrtype upper as.
3-52 Lower a~
The moa m u e n t orientation relationship associateci with the Bagaryatski
relationship [ 1 9501, is also often found in ae as: (00 1}, 11 { 2 1 1) and (1 00), 11 (07 1)
3.5 2 . 1 Spanos-Fanwkironson Model
Spanos et al. [1990] proposed a diffusion-controlled growth mode1 for lower ae based
on the TEM observations. As shown by sketches in Figure 3 5 3 , a single, fargely 8-fiee a
plate forms the "spine" of a lower as plate. "Secondary a plates" then fom predominantly at
one broad face of the "spine" by edge-to-face sympathetic nuchion, their broad faces lying at
about 5 5 to 66" with respect to the longitudinal axis of the "spine." 0. again precipitating
pnmarily from y at y/a boundaries wehemann rr al-. 1972b; Huang and Thomas, 19771, foms
at the broad faces of the "secondary a plates," often in the gaps between adjacent "secondary a
plates." [Spanos et al., 19901
Figure 3.5.3 Sketch of the difision-controlled rnechanism for lower a~ formation [Spanos et al., I W O ] .
3.5 -2.2 Bhadeshia Theow
Mer displacive formation of a plates in lower ae, Bhadeshia [1992] cunsidered excess
C atoms in the a plates is removed by two sirnuitanmus mechanisms: the precipitation of 0
within a or diaision of C into the residuai y. In lower ae, 8 precipitatioa dominates.
To explain the single variant of intralath 8 in lower ae, Bhadeshia [1990] reported
that Iower baùùtic a has a habit plane of approxkmte (0.76 1,O. 169, 0.626), and the intralath 8
precipitates on (1 Z), to form a single variant of 8 laying - 57' to the bainitic a ais. For
lower as containing both intra- and interlath 8, intraiath 8 precipitates first, followed by
the precipitation of interlath 8 in a way identical to that in upper as. The precipitation
reactions in lower a~ proposed by Bhadeshia are as follows:
in case 1, it is cunsidered that sutFcient C is tied up at the dislocations so that €4, a
transational 0, is not present [Kahsh and Cohen, 19701.
Ohmon [1989] postulated a displacive mode1 according to Bhadeshia's theory
(Figure 3 -5.4). Initidy, very smaii a subunits (about 200 x 30 x 50 nm) supersaturated with
C nucieate in a side-by-side fasbion by a displacive mechanism at the Cdepleted regions in the
vicinity of y grain boundaries. The cross section of the sublmit is probably in the shape of a
paralle10gram encloseci by both the habit p h e s and the lattice invariant shear plmes as in the
case of upper rn (Figure 3.5.44. The supersaturated C atoms in the a subunit will then be
rejected into the untratlsformed y and will buiid up at the interphase boundary. [f 0 platelets
nucleate epitmially on the lame invariant shear planes, which appear as ledges by the
coalescence of the subunits, in contact with both the a and the y, they form a row of 8 particles
and exhibit the Isaichev orientation relationship [Isaichev, 1947 with a (Figure 3.5.4b).
The Isaichev orientation relationship is quite close to the Bagaryatski relationship and
the two are f i cu l t to distinguish experimentaily @3hadeshia, 19921.
The repetition of these processes wifi form a typical lower as shown in Figure
3 -5.4~.
3.5.3 Thickening of ae
Rigsbee and Aaronson observeci in upper a~ that regularly spaced ledges and
dislocations of either edge or mked type aligned in a pardel fashion on the tenaces,
suggesting a thickening process by a ledge mechanism wgsbee and Aaronson, 19791.
Figure 3.5.4 Schematic representation of lower a~ growth [Ohmon, 19891. a) a subunit formation at an y grain boundary, b) (3 platelet nucleation on the side-surfaces of the individual
subunits, c) formation of lower a~ plate by the repetition of these processes
and the growth of 8 within the a plate.
On the displacive side, a~ thickening after lengthening is considered not possible
due to a loss of coherency of the interphase boudaries [Hehemann and Troiano, 19541,
and recent research on the cornpanson of sbon and long thne isothermal transformation
(IT) of a hi& Si-containhg steel supports such a conclusion [Tsuzaki et al., 19941.
Ohmon and Maki [ L W 11 maintained that the rearrangements of the dislocations
and step structures on the habit plane into more stable configurations occur via both
dislocation glide and the diffusion of substitutional atoms on it because the very coherent
&y lattices on the as habit plane interphase boundanes enable the driving force for the
thickening to be fairly srnail.
EFFECTS OF AUSTENITl3 CONDITION ON BALNITE TRANSFORMATIONS IN MICROALLOYED STEELS
4.1 DEFINITION OF BAINITE IN MICROALLOYED STEELS
ae microstructures obtained by controlied cooled d e r hot deformation in modem
rnicroaiioyed steels are usually very complex and there is some confusion in defining these
microstructures @Zdmonds and Cochrane, 1990; Araki and Enomoto, 19901.
4.1.1 Low-C Steels
A thorough definition of a* microstnictures in low-C microaüoyed steels is given
by the Bainite Cornmittee of the Iron and Steel Institute of Japan [Atlas of Bainitic
Microstructures, 19921. ae microstmctures are classified into major matmc phases: q,
that is fairly recovered, granuiar (lathless) bainitic ferritic structure with dislocated
substructures, and ma that is lathlike, 0-fiee bainitic a Iaths conseMng the prior y grain
boundaries. The minor secondary phases associated with bainitic structures are classified
into Bo or B2 (slightly C-enriched upper a~), Bu (C-e~ched upper a ~ ) and BL (higher C-
enriched lower cce) [Araki et al., 19921.
4.1.2 Medium-C Steels
a* types associated with medium-C steels are mostly lower a~ as the lower a~
usually forms in the temperature range between 400 O C and temperature [Ohmon and
Maki, 19911. Only in medium and hi&-C steels can M, be below 400 O C .
4.2 IUCRYSTALLIZED AUSTENITE
4.2.1 Bainite Nucleatioa
Yamamoto et a% [1995] reported tbat the density of intragranular micleation sites are
inaeased with the deaease in y grain size due to recrystalIization, but there are no &ea~ of y
grain s i x on growth, B. or microstructures. Bhadeshia [1982] commenteci that in any case,
unles site saturation ocnirs, as-start ùme should not Vary significantly with the usual range of
y grain sizes obtained foiiowiog commerd heat treatment.
In a 0.1C B-containhg steel, B was found to resegregate to the "fiesh"
recrystafiized y grain boundaries withui 1 second at 800-920 O C [Shikanai et al., 19881,
which suppresses the nucleation of aB.
4.2.2 Bainite Growth
While Umemoto et al. [1982] showed that the rate of CQ growth decreases with an
increase in grain size, Graham and Axon [1959] and Yamamoto et al. [1995] reported the
opposite, Le. there is no contribution of y grain boundary to the growth of bainitic a. It was
pointed out that the length of hths is decreased because of the refinement of y grains
[Yamamoto et al., 1 9951, but the bainitic subunit is not influenceci by either the y grain size or
the a~ lath size -en and Edmonds, 19781. Amoson [1969] found that a suffiCient
reduction in the y grain size can result in the replacement of the W~cbmtatten morphology of
a6 by that of grain-boundary dotriomorphs, but the volume hction of the traadomed
products depends on the temperature during continuous cooling aud does not chaoge with the
prior y grain Sze.
4.3 UNREKRYSTALLIZIED AUSTENITE
The energy is 1010 J/mol for a transformation [Ghosh and Oison, 19931, and - 400
J/moI for % ~ o m a t i o n phadeshia and Edmonds, 19801. Therefore, the effécts of the
stored energy in unrecrystauized y on tdormations is assumeci to be larger than on
transformations.
4.3.1 Bainite Nucleation
Deformation below Tm increases the density of nucleation sites and raises B. during
CCT, leading to an increase in a~ hardness and volume fiaction, which was enhanced by
increased cooling rate [Huang et al., 19931. Similady, transformation was reporteci to be
accelerated by deformahm below Tm during ïï Bwards and Kennon, 1974; Freiwillig et al.,
1 9761.
From cornparison of ae traasformations in a Nb-bearing and a Nb-fiee steei,
Yamamoto [1995] condudeci that the strain-induad Nb-rich precipitates did not infiuence the
0 1 ~ transformation.
4.3.2 Bainite Growth
It was reportai that heavy defiormation (i~e., 50%) below Tm remarkably reduces the
Iength of a l a h and the number of laths in the same orientation, r d h g in a sigdicant
refinement of lath Bjiwara et al., 19951. The nucleation of % on dislocation ce11
boundaries fomed in deformeci y grains WcQueeq 1977; McQueen and Jonas, 19751 was
proposed to account for this phenornenon mjiwara et al, 19951.
4.4 CRYSTALLOGRAPHY
Generally, the qstaUography of a~ is not changed due to i
recent results are given in Table 4.3.
Table 4.3 Variation of Bainite Crystaliography with TMP
'he
As-Reheated y Recrystallited y UnrecrystaIIized y Reference Y I ~ K-S K-S K-S Yamamoto et
ai.. 1995 habit plane (1 1 l), (1 il)a (45 1 }a Okaguchi, 1991
(i10}s{451), {110}=-{451}a lath adjacent laths adjacent laths most adjacent laths orientation have same have same have different Fujiwara et al.,
orientation orientation orientations 1995
For Bantainhg steels, B. was reporteci to be iowered [Hayashi et al., 19941 or hardly
changed Fujiwara, 19941. When Nb and B are added together, a synergistic effect of Nb
and B was found to promote the ae transfomation [Numg et al., 19931 due to: 1) Nb can
effectively retard recrystallization, and B has sufficient tirne to d w s e to the vicinity of y
grain boundaries to increase the hardenabiiity of y [Tamehiro et al., 19861, 2) Nb can
decrease the difisivity and activity of C in y, the dissolved Nb might therefore protect B
from forming BC pakasugi et al., 19811, and 3) solute Nb has itself a profound effect in
preventhg a kom forming makasugi et QL , 1 98 11.
B was found to resegregate to the unrecrystallized y grain boundaries and some
deformation bands just after deformation [Shikanai et al., 19881, which is beiieved to
apply a restraint to the nucleation of aB.
In summary, although there has been extensive study of a~ transformation
mechanisrns in steels, there is only limited understanding of the effects of austenite
condition on these processes. In view of the strong interest in therrnomechanical
processing of microaüoyed steels, this research was designeci to systematicdly study the
effects of austenite condition on bainite transformations.
5.1 MATERIALS
The materials used in this investigation were a 15B 13MA steel and a SAE1345M
steel, which were suppiied by commercial producers as hot-roiied bars with diameters of
12 mm and 27 mm, respectively. 1 SB 13MA is a B-alloyed low-C bainitic bar steel C'L"
steel), and SAE1345M is a Nb-containing medium-C forging steel ("M' steel). The
chernical compositions of the steels are given in Table 5.1 .
TabIe 5.1 Chernical Compositions of L and M Steels
5.2 THERMOMECHANICAL PROCESSING SIIMULATIONS
A Materials Measuring Corperation quench-deformation dilatometer was used to
perform various TMP treatments by varying deformation schedules and cooiing rates.
5.2.1 Dilatometer Set-Up
The cylindrical dilatometer samples of 4 mm in diameter by 8 mm in height were
machined with their height axis dong the longitudinal ( rohg) direction of the steel bars.
An "S" type Pet-lO%Rh thermocouple was spot-welded onto the sample surface to
control temperature to _+ 3 OC. Shims of O. 1 mm thick Mo sheet were also spot-welded to
each end of the sample to provide sorne lubrication dunng deformation.
Shown in Figure 5.2.1, the sample was set up between two Sic heads, and heated
by the induction coil in a vacuum of 10" torr. Deformation was performed by the
hydraulically driven Sic heads to a strain of 50% or 25%. Fast cooling rates of 30- 130
"Us were controlled by helium gas flowing out through the quench coil, slow cooling
rates of 0.01-2 " U s by the induction fimace itself. and intemediate cooling rates of 5-30
OC/s by helium gas and the fumace together. Change in sample length (AL) dunng phase
transformation caused a displacernent of the quartz rods, which was detected by the linear
voltage differential transducer (LVDT). The total range of the LVDT is + 2.5 mm with a
linear uncertainty of 0.25% full range [Collins and Barry, 19841. During continuous
cooling, AL,, time elapse (Ai) and temperature (7) were acquired by the controlling
computer. In addition, an X-Y chart recorder recorded cuves of AL vs. T at the
magnification of O. l pm/cm.
5.2.2 Design of TMP ScheduIes
The TMP scheduies, shown schematicaiiy in Figure 5.2.2, were designed as
follows.
5.2.2.1 Austenitization Temperature
To be compatible with indusuial practice, al1 samples were reheated at 5 " U s to
1 180 OC, and held at this austenitization temperature for 15 minutes.
5.2.2.2 Austenite Conditions
Three distinct y conditions were produced: 1) y as-reheated, 2) y deformed above
Tm (recrystailized y), and 3) y deformed below T,R (unrecrystallized y).
TNR was first estirnated using Equation 2.9 for the two steeis, then a series of
samples were defomed by 50% strain at temperatures around the estimated TNR, held for
10 seconds, and quenched to room temperature. The highest temperature below which a
100% recrystallized microstructure did not appear was taken as TYR.
5.2.2.3 Cooling Patterns
Samples were subjected to either continuous cooling or isothermd holding.
1) For the CCT expenments (CCT), dilatometric samples as-reheated were cooled
at 0.5-130 OC/s to room temperature (Figure 5.2.2a), or cooled at 2 "C/s to the
deformation temperature ( T d ) to receive a single compressive deformation of 50% strain
(Figure 5.2.2b), or double deformations of 25% strain each (Figure 5.2.2c), followed by
10 seconds hoIding and 1-130 *C/s cooling. CCT was aiso interrupted at various stages
by quenching at 130 OC/s to room temperature to obtain the panially transformecf
microstructures.
2) For the IT experiments (IT), dilatometnc samples were cooled at 100 OC/s from
the reheat temperature or Td to the desired IT temperature for 5-3600 seconds
transformation, followed by 130 OC/s quenching (Figure 5.2.2d).
Quartz Rods (to LVDT) Raten
Induction CMI L Quench Coil
Figure 5.2. I Dilatometer set-up. [Nelson, 19961
(c) (a Timc
Figure 5.2.2 Designed TMP schedules.
5.2.3 Determination of CCT Diagrams
During y CCT, the onset of a phase transformation was interpreted as the point on
the dI. vs. T curve where a deviation was identified (Figures 5.3. l a-e [Eldis, 19771). then
ail the determined transformation start (subscript "s") and finish (subscript "f')
temperatures were ploned on cooling curves on a T vs. log t (time) diagram or CCT
diagram (Figure 5.3. If). To compare aii the CCT diagrams on the same time base. the
cooling stan temperature was taken to be 750 OC. This temperature is between the lowest
Td of 780 O C and the highest Ar3 of 720 OC.
To determine the transformation start and finish temperatures, one method is to
find the intersection of tangents to adjacent segments of the cooling curve (Figure 5.3.2a);
another method is to find the point of initial deviation from the cooling curve, al1 having
an accuracy of & 10 O C (Figure 5.3.2b). The former method leads to a measurement - 10
OC lower in transformation start temperatures and - 10 OC higher in finish temperatures
than the latter method. During CCT, a deviation nom the AL us. T cuve occurs only
when a certain amount (- 5%) of volume changes has ocnirred [Collins and Barry, 19841,
so the detected transformation start temperature is lower than the actual start temperature,
hence the latter method was used in this study.
5.3 MICROSTRUCTURAL CHARACTERlZATION
During deformation, the cylindrical sarnples usuaily barre1 because of friction
between the sample and the platens, so the Iargest strain occurs in the middle of the
sample and gradually decreases towards the ends [Le Floc'h et al., 19871. Therefore, a
dilatornetric sample was sectioned parailel to its longitudinal axis or compression direction
using a diamond cutter to make specimens for microstructural characterization. The
region near the sample centre was exarnined using opticai, scanning electron microscopy
(SEM) and transmission electron microscopy (TEM).
5.3.1 Optical and Scanning ELectron Microscopy
Microstructures of specirnens polished with 0.5 pm alumina powder and etched
with 2% Nital were charactenzed by both opticai microscopy and a IEOL JSM-840 SEM
microscope operated at an accelerating voltage of 10 kV and a working distance of 15-39
mm. An EDS (Energy Dispersive Spectroscopy) X-ray detector was used with spot
scanning mode to identify the types of coarse particles. y grain boundaries were reveded
by examining the deformed+quenched sarnples described in 5.2.1.2 using a saturated picric
acid+wetting agent solution maintained at 70 O C . The rnean y grain diameter was
measured by a linear intercept method, and the volume fraction of each microstmcturai
phase by the point counting technique on SEM rnicrographs using a 10 by 10 grid [Metals
Handbook, 19851.
5.3.2 Transmission Electron Microscopy
A Philips CM 20 transmission electron microscope (TEM) was used at 200 kV to
characterize: 1) the precipitates in y by C extraction replicas, and 2) the microstnictural
details by thin foils.
5 -3 -2.1 S~ecimen Pre~aration
C replicas were prepared as follows: first, a thin C layer was deposited onto the
surface of a lightly 2% Nital-etched sample using a JEOL evaporator in a vacuum of about
10'~ torr. A dark brown colour indicated a suitable thickness of C layer that was scored
into - 2 mm by 2 mm squares using a scaipei. Lastly, the C squares were Boated off in 10
% Nital, washed by distilled water, mounted on the 3 mm diameter copper gnds and drkd
on filter paper for examination.
Figure 5.3.1 Schematic illustration of typical dilatometer records on continuous cooling. (a-e) AL vs. T records for various cooling cycles; ( f ) cooling cycles for (a- e). subscript "s" is start temperature, is finish temperature, T is temperature, L is length, r is time [Eldis, 19771.
Figure 5.3.2 Two methods of determining the transformation start and finish points.
To make thin foiis, 0.1 mm thick slices were cut dong the direction parallel to the
compression axis near the middle region of the dilatometer samples. The slices were
ground using a specially made sample-holder to - 80 pm in thickness, and punched into
discs of 3 mm diameter. The discs were electropolished using a twin jet polisher operated
at 25 V and -10 OC in an electrolyte of 10% perchloric acid+75% acetic acid+l5%
methanol until perforation occurred. To preserve precipitates, a &Os-based electrolyte
polishing solution was used.
5 -3.2.2 TEM Examination
In examining replicas, a Si(Li) detector with a minimum probe sue of 7.5 m and a
Noran (Voyager) rnicroanalysis system were used to carry out the chernical analysis of
precipiiates. Precipitate size was measured directly on the TEM micrographs. The
irregular morphologies were converted into spheroids of equivalent volume. The number
of particles per unit a r a was utilized to descnbe the particle density. Usually 100- 1000
particles were measured for each size distribution.
The crystdlographic orientation relationship (COR) of the associated phases were
determined using composite diffkaction patterns and standard stereographic analysis. The
incident beam direction of each phase was identified separately, and dl were transferred to
a (00 1) stereograph. COR was determined by those superimposed stereographs using the
rotation maneuvers described by Edington [198 11.
5.4 HARDNESS TESTING
A Vickers Microhardness tester was used to determine the hardness of the
dilatometer samples. A load of 30 kg was used, and four indentations were made for each
hardness measurement.
6.1 AUSTENITE CONDITIONS
From the quenched samples representbg the various y conditions (see 5.2.1.2), the
diarneters o f the recrystallized or unrecrystaiiized y grains, and the types and distributions
of precipitates in y were examined.
6.1.1 Evolution of Austenite Grain Structure
Using the chemicai compositions in Table 5. L and Equation 2.9, T N ~ temperatures
were calculated to be 1233 O C and 1276 OC for L and M steels, respectively. By the
deformation+quench experiment, TM temperatures for 1180 OC-reheated L and M steels
were detennined to be 950 OC and 1025 OC, respectively. The discrepancy between the
rneasured and calculated Tm values was rnainly due to the term (6645AB- 644JNb) in
Equation 2.9.
6.1 . 1 . 1 Austenite Grain S tnictures
The different y grain structures produced by the TMP treatments (given in Figure
5.2.2) for both L and M steels are summarized in Table 6.1. Generally, the y grain size
(Dy) decreased with decreasing Td. L steel as-reheated started with Dr of 21 Pm.
Deformation (50% strain if not othenvise indicated) above Tm at T d of 1000 O C refined Dy
by 20% to 17 pm through recrystallization. Defomation below T , at 850 or 780 OC
produced a "pancaked" y grain structure with a width of 15 or 14 pm, - 33% smailer than
D, for L steel as-reheated. M steel as-reheated had a Dy of 45 pm. Deformation above
TNR decreased D, by 42% to 26 pm for Td of 1080 O C , or by 65% to 16 pm for T d o f
1025 O C . Defomation below Tm reduced the as-reheated Dy by - 70% to 14 (Td = 950
O C ) or 13 Pm. (Td = 850 OC). It is obvious that the decrease in 4 by TMP is more
significant in M steel than in L steel.
Table 6.1 Surnmary of Austenite Grain Structures
6.1.2 Evolution of Precipitate Distribution in Austenite
r
Steel
L
M
6.1.2.1 L Steel
Distributions and morphologies of precipitates present in the various y conditions
are illustrateci by the typical TEM replica micrographs in Figure 6.1.1. The results of the
* width of y grains
&, O c
as-reheated 1000 850 780
as-reheated 1080 1025 950 850
4 Pm 2 1 17 15. 14.
45 26 16 14. 13.
Grain S tmcture recrystallized recrystdlized
pancaked pancaked
recrystallized recrys tallized recrystallized
pancaked pancaked
number density (p) and the corresponding average diameter (4 of precipitates are given in
Table 6.2. Due to the extraction efficiency, the measured p in Table 6.2 is only a relative
representation of precipitate number density, which should be lower than the actual
number density. The precipitate size distributions under each y condition are s h o w by the
histograms in Figure 6.1.2. It is noted that:
1) In L steel as-reheated, only a Iow p of spheroidai (Nb,Ti)-nch precipitates of - 15
nm in diameter were present (Figures 6.1.1 a and 6.1.2a). Severai large MnS particles up to 2
pm were also found.
2) L steel deformed above TM at 1000 O C contained a high density of strain-
induced precipitates. By the size distribution, there were a group of fine precipitates of - 15 MI in diameter, and a group of medium-sized precipitates of - 35 nm (Figures 6.1.1 b
and 6.1 -2b). The average precipitate diameter was larger than that of L steel as-reheated.
Some (Nb,Ti)-rich precipitates were observed to f o m at a pre-existing coarse MnS
particle. Generally on the EDS spectra, the coarse (Ti,Nb)-nch particle contained more Ti
than Nb while the fine (l'%,Ti)-nch particle had more Nb than Ti; the intensity peaks of C
and Cu were £?om the C replicas and the Cu grids nipporting replicas, respectively (Figure
6.1.3).
3) L steel deformed below TM at 850 O C or 780 O C contained a low number
density of fine (Nb,Ti)-rich precipitates (Figure 6.1. Ic and Table 6.2) , - 3 5 nrn in average
diameter (Figures 6.1 .2~ and d). A similar precipitate distribution was observed in L steel
deformed above TNR at 780 O C at 25% strain and below TnR at 780 OC at 25% strain
(Figure 6.1 -2e).
Figure 6.1.1 TEM replica micrographs showing precipitate distributions in deformed+quenched L steel. a) as-reheated, b) Td = 1 O00 OC, C) Td= 780 OC.
Table 6.2 Number Density @) and Average Diameter (d) of Precipitates
5 15 25 35 45 55 65 75 85 95
Partide Diameter, nrn
1 ~ V L T
r d cc)/& (%) as-reheated
1080/50 -- -- 26477 in?c/cn
25k62
M Steel L Steel p, 1/m2
66 p, hm2
327 4
1 08k50
4 3 1 S 3
5 1 5 Z 3 5 4 5 5 5 G 7 5 & %
Parücle Diameter, nrn
Particle Diameter, nm
Figure 6.1.2 Precipitate size distributions in L steel. a) as-reheated, b) Td = 1 O00 OC, C) Td = 850 OC. d) Td = 780 OC. e) 7'' = 1000 "C (25%~) and 780 OC ( 2 5 % ~ )
120
1101
1 001
901
do(
cn u 70t
8 S "'"
soa
400
300
ZOO
100
O
l / l l l
1 lui
r noi
901
dom
7uu
600
500
4 on
300
m ZOO
Energy (keV)
4
Figure 6.1.3 @&,Ti)-rich particles nucleated on a MnS particle in L Steel deformed at 1000 OC. a) TEM bright field, b) TEM dark field using
reflection of the (Nb,Ti)- rich particle,
c) EDS spectrum for MoS, d) EDS spectrum for coarse
(Ti,Nb)-rich particle, e) EDS spectnim for fine
@&,Ti)-rich particle.
Energy &eV)
6.1.2.2 M Steel
Distributions and morphologies of precipitates present in y deformed above and
below T , are illustrated by the TEM replica micrographs in Figure 6.1.4. By EDS
analysis, most precipitates were (Nb,V,Ti)-rich and (Nb,V)-ich, while some were Nb-rich
and VC. For each y condition, the number density and the corresponding mean diameter
of precipitates are given in Table 6.2, and the precipitate size distributions are shown by
the histograms in Figure 6.1.5. No meaningfid statistics were made for M steel as-
reheated because oniy very few spheroidal or irregular 100- 1000 nm (Nb,Ti)-nch particles
were present. However afler deformation, fine strain-induced precipitates constituted 70-
90% of al1 the precipitates. It is noted that:
1) M steel deformed above Tm at 1080 O C contained fine (- 15 nm)
sphericakubidfaceted strain-induced precipitates rich in Nb, Nb+V, or Nb+V+Ti and V,
which were aligned in lines (Figures 6.1.4a and 6.1 Sa). h M steel deformed above TNR
at 1025 OC, the strain-induced precipitates were extra fine (- 5 nm) in average (Figure
6.1.5b). In some areas, precipitates were present in clusters with new precipitates
nucleating at the pre-existing ones. For example, Figure 6.1.6 showed that three V-nch
precipitates forxned on a coarse (Nb,Ti)-rich precipitate.
Figure 6.1.4 TEM replica micrographs showing distributions of strain-induced precipitates in deformed+quenched M steel. a) Td = 1080 and b) Td = 850 OC.
5 15 25 35 45 55 65 75 85 95
Particle Diameter, nm
5 15 25 35 45 55 65 75 85 95
Particle Diameter, nm
5 15 25 35 45 55 65 75 85 95
Particle Diameter, nm
5 1 5 2 5 3 S & S E E S %
Particle Diameter, nm
Figure 6.1.5 Precipitate size distributions in M steel. a) Td = 1080 O C , b) Td = 1025 O C ,
C ) Td = 950 O C , d ) Td = 850 O C ,
e ) Td = 1080 O C ( 2 5 % ~ ) and 850 O C (25%~).
Figure 6.1.6 niree V-rich particles nucleated on a coarse (Nb,V)-rich particle in M steel deforrned at 1025 OC. a) TEM bnght field, b) EDS spectnun of V-rich
particle, C) EDS of (Nb,V)-rich
particle.
2) In M steel deformed below Tm at 950 OC, the strain-induced precipitates were a
high density of extra fine (- 3 nm) precipitates and some medium-sized precipitates (- 25
nm), al1 being nch in Nb+V+Ti or M+V (Figure 6.1 Sc). M steel deformed below TNR at
850 OC had Iarger strain-induced spherical precipitates in a lower number density: fine
ones were 8- 15 nm in diameter distributed in the matrix and medium ones were 60- 1 50 nm
distributed dong curved lines which define areas - 1 Pm in size (Figures 6.1.4b and
6.1 Sd). In M steel deformed 25% above Tm at 1080 O C and 25% below TNR at 850 OC,
very coarse undissolved precipitates (up to 3 prn) CO-existed with a high density of strain-
induced fine precipitates (- 5- 15 nm) (Figure 6.1 Se). Pure VC precipitates were also
detected.
6.2 BABVITE TRANSFORMATION KINETICS
6.2.1 Microstructure-Based Definitions of Bainite
The primary transformation products of y are a, P, a~ and various types of a ~ .
After CCT, final microstructures are nonnally various combinations of these
transformation products and, thus, are very complex.
This study considers a ~ , the phase continuously transfomed in the intermediate
temperature range between that of a/P and a ~ , and classifies a* types according to their
bainitic a morphologies, their formation temperatures and distributions of C-rich phases
such as iron carbides (8) or retained y (y,). Two morphologies of bainitic cr are identified,
lathlike (aeL) and lathless (aeh). Following Ohmori [1971], a B L is fûrther classified into
BI (aB laths + interlath a&~), BI( (aB la& + interlath e), Bm ( a ~ laths + intralath 8). A
fourth type of aeL was identified in this study, BN (aB laths + interiath aM/y~/B + intralath
0). BrBN dl have a lathlike rnorphology, and it will be shown in Sections 7.3 and 7.4 that
they also have sirnilar transformation mechanisrns. Accordingly, a B h is îürther classified
into B: (granular a~ + aM/yR islands) and B: (elongated a~ laths + interlath a&~) . The
subscript "I" is used for B~~ and B: because of their 8-fïee feature. It wiil be shown that
the formation temperature decreases in the order BI~-B;-BI-BN- The definitions of the
various a~ constituents are summarized in Table 6.3.
Table 6.3 Definitions of Continuously Transfomed Bainite Types
The various ae types are iiIustrated by the micrographs in Figures 6.2.1 and 6.2.2.
Under SEM, B: was present as granular as subgrains with equiaxed islands; B:
had the elongated ae subgrains with corne inter-mbgrain aM/yR; Br and Bn appeared as
dark a~ laths with slightly grayish or white interlath films of yR (rnost y~ in Bi transforms
to aM) and 9, respectively; intralath 0 in Bm stood out as white dots in secondary electron
SEM images; Bw showed a characteristic lenticular morphology. Under TEM, the details
of these a~ morphologies and the complex distributions of C - e ~ c h e d phases (8, a~ and
yR) were clearly clarified.
interlath a d y ,
6.2.2 CCT Behaviour of L Steel
The CCT diagrams for each y condition are show in Figure 6.2.3. The underlined
numbers represent the cooling rate in 'Ch, and the parenthesized numbers are the
Vickers ' hardness of the final microstructures.
interlath 9 intraiath 0 interlath u ~ / Y ~ / ~ + htraiath 8 -
+ aM/yR + ~ M / Y R
Figure 6.2.1 a~ morphologies-SEM: a) Br and Brrr, b) Brr and Brv, C) B~~ and B:.
Figure 6.2.2 a~ morphologies-TEM: a) Br and BUI, b) Bu Biv, c) BI?
10 1 O0
Time, s
- -- - -. _ ---. _ - - -.- . - -- ;zr= Austenite Grain Size =20 micron
10 1 O0
Time, s
Figure 6.2.3 CCT diagrams of L steel. a) as-reheated, b) Td = 1 000 OC, C) T' = 850 O C , d) Td = 780 O C , e) Td = 1 O00 OC (25%~) & Td = 780 O C (25%~).
Figure 6.2.4 AL- T curves at typical cooling rates of a) 130, b) 5, and c) 1 OC/s.
The TMP scbedules are given in Figure 5.2.2. AL-T curves obtained at typical
cooling rates of L steel are shown in Figure 6.2.4, on which the start andor finish
temperatures of a, as and a~ are represented as Fs, BJBf and M m f y respectively. At the
highest coolhg rate 130 "Ch, y decomposed at 425 OC to produce a final microstnicture
of mostly a~ and some asL, but it was not possible to distinguish B. fiom M. on the ALJ
curves, hence
were cleariy
425 OC was taken as both B. and M, (Figure 6.2.4a). At 5 T/s, B. and Br
identified as 590 OC and 425 OC, respectively, as a complete ae
microstructure was formed (Figure 62-46). At 1 'Us, B, was 6 15 O C , preceded by the
formation of a at Fs of 660 OC (Figure 6.2.4~).
Table 6.4 lists the volume fraction of each phase present in the final microstructure
obtained at various cooling rates in the as region.
Table 6.4 Volume Fraction of Transformation Products in L steel
-*-- TMP Treatment -----
TA. O c COOL "C/S
as-reheated 10 1 O00 10 850 10 780 10
1 000&780 10
as-reheated 5 1 O00 5 850 5 780 5
1 000&780 5
as-reheated 1 1 O00 1 850 I 780 1
1 OOO&78O
Volume Fraction. %
6.2.2.1 L Steel As-Reheated
Figure 6.2.3a shows that a~ was obtained over a wide range of cooling rate, 0.5-
130 "Ch, in the temperature range 6 15-3 00 OC, while or and P transformations occurred at
cooling rates below 1 O C k aeL transformed at lower temperatures and higher cooling
rates, 130- 10 'Ch. a** appeared at higher temperatures and slow cooling rates, 5- 1 OC/s.
The hardness of the final structures decreased with the descending cooling rate.
6.2.2.2 L Steel Deformed Above Tm
a~ appeared in the kinetic range 130-1 "Cls and 580-300 OC (Figure 6.2.3b).
Compared with L steel as-reheated, L steel defomed at 1000 OC had a CCT diagram on
which the as region was as though "squeezeci" fiom the top and right side. Specificaily,
B. was lowered by 15-40 OC at 30- 1 OC/s causing a decrease in a~ volume fraction. For
example, a* volume fraction was decreased by 50% at 1 OC/s (Table 6.2). On the other
hand, Fs and P, were raised by 10-40 OC . The a "nose" cooling rate below which a starts
to form was increased fiorn - 1 "Cls to - 5 OUs. The hardness was higher at 130-10 "C/s
( a B L + a M range) but lower at 5- 1 "Ch (a$ range).
6.2.2.3 L Steel Deformed Beiow TNR
The a* kinetic range was 130- 1 ' U s and 620-280 O C for 7'' of 850 OC, and 130- 1
OC/s and 6 10-250 OC for T d of 780 OC (Figures 6.2.3 c and d). Cornpared with that of L
steel as-reheated, B. or M, decreased by 25-50 OC at high cooling rates of 130-30 'Us,
but ae transformation was significantly accelerated at 30-5 OC/s as the aBA nose cooling
rate increased fiom 5 OC/s to 30 " U s , resulting in an increased a e A but decreased a B L . So
the tota! ua volume fraction kept relatively constant. At slow cooling rate of L 'Us, B,
was constant but 70%-85% a e A was replaced by a with a nose rate being accelerated to
above 10 "Cfs in contrast to 1 "C/s of L steel as-reheated. The hardness of the final
microstmctures decreased with decreasing cooling rate.
6.2.2.4 L Steel Deformed Above and Below Tm
The a~ kinetic range was 130-1 OC/s and 600-200 OC for 25% strain at 1000 and
780 OC (Figure 6.2.3e). Compared with that of L steel as-reheated, the a~ kinetic range
was not significantly altered by the double deformation. However, unlike L steel
defonned at 780 OC, the a nose cooling rate was only slightly increased in this case. The
hardness of the final microstructure was slightly lower than that of L steel as-reheated and
cooled at the sarne cooling rate.
6.2.3 CCT Behaviour of M Steel
The CCT diagrams for each y condition are illustrated in Figure 6.2.5, and the
typicd AL-T curves are shown in Figure 6.2.6. The volume fractions of each phase in the
final microstmcture obtained at representative cooling rates are given in Table 6.5. In M
steel, ae exhibited a lathlike morphology (aeL) and a had an allotriomorphic morphology,
fiequently associated with y grain boundarîes (GBa).
6.2.3.1 M Steel As-Reheated
Cornpared with the a~ kinetic range spanning 130-0.5 OCls and 6 15-300 OC in L
steel as-reheated, the corresponding range for M steel as-reheated was 20-1 " U s and 520-
400 OC. Therefore B. and Br of M steel were - 100 OC lower and higher than that of L
steel, respectively (Figure 6.2. Sa).
Table 6.5 Volume Fraction of Transformation Products in M steel
TMP Treatment .*.-*~.*.-----.-.-*.~*---.-UIIUIUIIUIUIIUIUIIUIUIIUI~~
as-reheated 10 1080 10 1 025 10 950 10 850 10
1080&850 10
as-reheated 5 1080 5 1025 5 950 5 1 850 5
1 1080&850 5
as-reheated I 1 080 1 IO25 1 950 1 850 1
1 080&850 1
VoIume Fraction, % .- UIIUI*--.-..-...--.l-.~-.--.-.----.*--~-*---.~-~-.--**.---
a P a~
From 20 to 1 OCls, B, gradually increased fkom 400 to 520 O C , resulting in a small
increase of aeL volume fraction fiom 3 1% to 44%, and a decrease of a~ from 67% to
zero. The hardness of final microstmctures decreased with decreasing cooling rate.
The hardenability in terms of the ability to obtain a~ for M steel was obviously
higher than L steel as complete a~ microstmctures were obtained at cooling rates of 130-
20 OC/s. The a~ region was completely separated from the aa region, so a non-
transformation region where y was metastab le exist ed before transforming into U M (Figure
6.2.6a). As y transformed continuously, it was not possible to distinguish Fs and P, h m
B. (Figure 6.2.6b).
6.2.3.2 M Steel Deformed Above Tm
Generally, the ae kinetic ranges were shrunken compared with M steel as-reheated
(Figures 6.2.5b and c). M steel deformed at 1080 OC had an ae kinetic range of 10-1 "Cls
and 490-370 OC. Bs decreased by 40-60 O C , but the aa volume fraction oniy decreased by
less than 10%; similarly, M steel deformed at 1025 OC had an a~ kinetic range of 10- 1
"Ch and 460-360 OC. The ae volume fraction was not decreased although Br decreased
by 20 OC. For both T,, ag was completely replaced by 70% P and 30% a at 1 " U s . The
a nose cooling rate was increased corn 10 OCfs for M steel as-reheated to above 30 O C / s
and F, was raised by 60- 100 OC.
The hardness of the final microstructures decreased with decreasing cooling rate.
The hardness was higher in M steel deformed above Tm than in M steel as-reheated at the
higher cooling rate, 130- 10 "Ch, but becarne equal at slower cooling rates.
-. - - * - - &en& Grain Size = 45 micron
Time, s
Aostenite Grain Size = 26 micron
* L W..
F
Time, s
Austenite Grain Size = 16 mivon
10 100
Time, s
- --y_ - - - Austenite Grain Sire = 19 micron
Figure 6.2.5 CCT diagrams of M steel: a) as-reheated, b) Td = 1080 OC, C) T' = 1025 OC, d) Td = 950 OC, e) Td = 850 OC. fi rd = 1080 O C (25%~) & = 850 OC (25%~).
B, (590 O C )
b)
Figure 6.2.6 AL-Tcurves at typical cooling rates oE a) 5 and b) I "Cls.
6.2.3.3 M Steel Defonned Below Tm
Compared with M steel as-reheated, M steel deformed at 950 O C had a raised B, at
20-5 "Us, but a relatively stable a~ volume fraction. M steel deformed at 850 OC had an
unaltered a~ kinetic range and a slightly decreased a~ volume fraction (Figures 6.2.5d and
e). At 1 OCIs, only a and P were present.
Meanwhile, the a nose cooling rate of a was increased from 10 "Cls to above 30
"Ch, and Fs was raised by - 60 OC. M s was significantly lowered at the higher cooling rate
range, 1 3 0- 1 0 O C/s.
Compared with that of M steel as-reheated, the hardness of final microstructures
of M steel deformed below T N ~ was higher.
6.2.3.4 M Steel Deformed Above and Below Tm
Compared with that of M steel as-reheated, the as region in this case was oniy
slightly altered with a slight decrease of the a~ volume fiaction. Unlike the single 50%
deformation, the double deformation did not "close up" the a* region at 1 "Cls, instead,
34% a~ remained (Figure 6 . 2 3 ) .
Meanwhile, Fs was significantly promoted, and a transformation was significantly
accelerated as the a nose cooling rate was increased to 130 "Ch.
6.3 MICROSTRUCTURES OBSERVED BY SCANNING ELECTRON MICROSCOPY
Three groups of experiments were cmied out: 1) CCT, wbich determined the
effects of y condition and cooling rate; 2) IT, which defined the temperature ranges for the
various a~ transformations; and 3) interrupted CCT and IT (partially transformed), which
identified a~ nucleation sites.
In this section, generai features of microstructures resulting from the above
experiments and characterized by SEM are described. The details of the rnicrostmctures
as determined by TEM are described in Section 6.4.
6.3.1 Austenite As-Reheated
6.3.1.1 L Steel
The SEM microstructures produced by CCT (referred to as "SEM CCT
microstnictures") of L steei as-reheated are shown in Figure 6.3.1. At high cooling rate of
130 "Ch, the microstructure was a mixture of au (major constituent) and some lathiike as
(aeL). In addition to focming at prior y grain boundaries, aeL (mainly BI) was also
obsenred to nucleate at an annealing twin boundary and grow into one side of the
anneaiing twin, delineating the twin boundary as a distinct straight line (arrow in Figure
6.3. la).
Figure 6.3.1 SEM CCT microstructures - L steel as-reheated and cooled at: a) 130, b) 30, and c ) 10 OC/s.
At 30 or 10 OC/s, aBL formed sympathetically at pnor y grain boundaries, and grew
into bundles, thereby delineating the pnor y grain boundaries. In coarse y grains, the new
a B L laths were observed to fom at pre-existing aBL laths. At 1 0 OC/s, - 5 % BI was
present (arrow in Figure 6.3. lc). Br had an average lath width of - 1.5 Pm. Bm bad a
lath width of - 0.8 Pm.
It is noted that aeL bundles propagated until impinging with other asL bundles, y
grain boundaries or twin boundaries, and each y grain usuaily contained more than one aaL
bundle.
At 1 OC/s, the microstructure was lathiess aa (aB9. Figure 6.2. (c shows both
granuiar and elongated a ~ / y ~ islands in BI and B ~ ~ , respectively. The y grain boundaries
in the B: region were weakly visible, and those in the BI^ region were not visible. B: had
an average a~ lath width of - 3 Pm.
The SEM microstnictures produced by IT (SEM IT microstmctures) of L steel as-
reheated are shown in Figure 6.3.2. For an IT time (tn) of 3600 seconds at 700 O C , the
microstructure was a + a~ (Figure 6.3.2a). .For In of 1800 seconds at 600 O C , the
microstructure was coarse aeA (BF) and some a, and the y grain boundaries were not
visible (Figure 6.3.2b). For ln of 3600 seconds at 500 OC, well-developed aaL and some
GB-nucleated P were present, so the y grain boundaries were distinct (Figure 6 .3 .2~ ) .
b)
Figure 6.3.2 SEM IT microstructures - L steel as-reheated and held at: a) 700 OC, 3600 seconds, b) 600 O C , 1800 seconds, c) 500 OC, 3600 seconds.
6.3.1.2 M Steel
The SEM CCT microstructures of M steel as-reheated are shown in Figure 6.3.3.
At 10 OCls or 5 'Us, the microstructure was ae in an a~ rnatrix with the CYB laths showing
the same variant as that of the underlying a~ 1athsAaths (Figures 6.3.3a and b). aBL had
two morphologies: the GB-nucleated lathlike at, bundles (Figure 6.3.3 dl), and the
intragranularly nucleated lenticular a~ laths (arrow in Figure 6.3.3b2). The former was
mostly BrBm bundles, the latter was mostly individual Brv laths. At 1 "C/s, GBa and P
were present at pnor y grain boundaries, and ae (mostly Bu) formed only in the middle of
the y grains. Some coarse Bn laths divided the aa region into smdler ones, limiting the
growth of other as (Figure 6.3 -3c).
The two intragranular nucleation sites of as identified in M steel as-reheated were
pre-existing u~ (arrow in Figure 6.3.3b2) and large Nb-nch precipitates (Figure 6.3.4).
The SEM IT rnicrostmctures of M steel as-reheated are shown in Figure 6.3.5 (ln
= 1800 seconds). Figure 6.3.5a shows the microstructure of a+P at 600 OC. At 500 O C ,
the formation of GBa was avoided due to 50 0(5/s-cooling fiom 1 180 O C to 500 OC, and a
complete ae (mostly Bu) microstructure was produced (Figure 6.3.5b). At 400 OC, the
microstructure was some a~ and mostly Bm containing coane 8. The pnor y grain
boundaries were not visible (Figure 6.3. Sc).
Figure 6.3.3 SEM CCT microstructures - M steel as-reheated and cooled at: a) 10, b) 5 , and c ) 1 'Us.
Energy
Figure 6.3.4 SEM CCT microstructures - M steel as-reheated. a) a~ nucleated at Nb-rich precipitate, b) The correspondhg EDS spectnim for Nb-rich particle.
6.3.2 Austenite Deformed Above Tm
6.3.2.1 L Steel
The SEM CCT microstnictures of L steel deformed above Tm are shown in
Figure 6.3.6.
Compared with L steel as-reheated (di the cornparisons are made between samples cooled
at the same cooling rate if not otherwise indicated), L steel deformed above Tm
contained more htragranularly-nucleated a~ laths. For example at 30 "Ch, a
predominantly intragranular nucleation of aB was seen, making the y grain boundaries very
faint (Figure 6.3.6a). At 10 "Ck, some intragranularly-nucleated a* was still seen, but as
formed mainly at y grain boundaries, delineating the prior y grain boudaries (Figure
6.3.6b). However, the morphology and the volume %action of a~ were essentialiy
unchanged in this coohg rate range (30- 10 OC/s).
Figure 6.3.5 SEM IT microstnictures - M steel as-reheated and held for 1800 seconds at: a) 600, b) 500, and c ) 400 O C .
The intragranular nucleation sites in for ae L steel were identitied as pre-existing
a~ laths, twin boundaries (e.g., weak traces of twin boundaries were present in Figure
6.3.6.al), and possibly precipitates because of arrays of radiaily distributed ae laths
(Figure 6.3.6a2).
At 5 'ch, the microstructure was mostly a e A and some a B L , and y grain boundaries
were not visible (Figure 6 . 3 . 6 ~ ) . At 1 OC/s, the microstructure contained a, some ae" and
P. Note that the aM/yR phase obtained at 1 'Us was coarser and more granular than at 5
OCIs (Figure 6.3 -6d).
The SEM intempted CCT microstructures of L steel deformed above TNR are
shown in Figure 6.3.7. When 10 OC/s-cooled L steel was interrupted at 500 OC by
quenching, a microstructure of rnainly a~ and some aBL was produced. Note that the
auto-tempered a~ laths were coarse and heavily etched because of the multi-variant
intralath 8 (Figure 6.3.7a). Cooling at 10 OCIs to 550 OC avoided the formation of a and
thereby the following 1 OC/s produced a distinct coarse aM/yR-containing B: al3 ,
structure nucleating at y grain boundaries (Figure 6.3.7b). When 5 "Ch-cooled L steel
was intempted at 520 OC by quenching, the microstructure was BI^, a~ and some BE. So
B; formed both fiom y grain boundaries and intragranularly above 520 OC, preceding the
formation of B~~ (Figure 6.3 -7c).
The SEM IT microstmcture of L steel deformed above TNR in Figure 6.3.8 (llT =
1800 seconds at 400 OC) shows both a~ and fine aeL (Bm) that contained very coarse
intralath 0 and delineated the y grain boundaries. The microstructure of 500 or 600 O C
resembled that in L steel as-reheated and is not shown here.
6.3.2.2 M Steel
The SEM CCT microstructures of M steel deforrned above TvR are show in
Figures 6.3.9 and 6-3-10.
In the case of deformation at 1080 OC and cooling at 10 'Us, a~ nucleated at y
grain boundaries and interiors without 8 precipitation (Figure 6.3.9a). At 5 "Ch, in
cornparison with that in M steel as-reheated, the amount of intragranular nucleated aeL
was decreased, but more GB-nucleated a B L bundles were present (Figure 6.3 -9b). At I
OC/s, the microstructure was totally GBa and P (Figure 6.3.9~).
The microstructures for Td of 1025 and 1080 OC were the sarne. As indicated by
the arrow in Figure 6.3.10a, an aeL lath laid across the y grain, dividing the y grain into
two "effective" grains, and the growth of other intragranular as was confined to those two
effective grains. At 5 'Ch, more ae and some P were present (Figure 6.3. lob).
An SEM 520 OC-interrupted CCT microstructure of M steel deformed at 1080 OC
and coofed at 5 "Cls is show in Figure 6.3.11, in which fine cle just began to form at the
grain boundaries.
Figure 6.3.6 SEM CCT rnicrostnictures - L steel deformed at 1000 OC and cooled at: a) 30, b) 10, c) 5 , and d) 1 'Us.
Figure 6.3.7 SEM interrupted CCT microstructures - L steel defonned at 1 O00 OC and: a) 10 OC/s to 500 OC,
quench, b) 10 OC/s to 550 O C ,
1 "C/s cool, and c) 5 OC/s to 520 O C ,
quench.
Figure 6.3.8 SEM IT rnicrostnicture - L steel deformed at 1 O00 OC and held for 1800 seconds at 400 OC.
Figure 6.3.9 SEM CCT microstructures - M steel deformed at 1080 OC and cooled at: a) 10, b) 5, and c ) 1 "C/s.
Figure 6.3.10 SEM CCT microstructures - M steel deformed at 1025 OC and cooled at: a) 10 and b) 5 OC/s.
Figure 6.3.11 SEM interrupted CCT microstructure - M steel deformed at 1080 OC and cooled at 5 OC/s to 520 OC, quench.
Figure 6.3.12 SEM IT microstructures - M steel defomed at 1080 OC and held at: a) 500 OC, 1800 seconds, b) 500 OC, 5 seconds, c) 400 OC, 1800 seconds, d) 400 O C , 5 seconds.
The SEM IT ( t f l = 1800 seconds) and interrupted IT (rrr = 5 seconds)
microstmctures of M steel deformed at 1080 OC are shown in Figure 6.3.12. At 500 O C ,
1800 seconds holding produced rnainly GB-nucleated (Figure 6.3.12a), while 5
seconds holding just saw a start of a~ nucleation at triple points of the y grain
boundaries, grain boundaries, and twin boundaries (Figure 6.3.12b). At 400 OC, 1800
seconds holding produced a very messy microstructure containing very coarse 8 (Figure
6.3.12~). By contrast, 5 seconds holding preserved the y grain boundaries and the lath
boundaries. Finer 8 was present in Bur (Figure 6.3.12d).
6.3.3 Austenite Deformed Below TNR
6.3.3.1 L Steel
The SEM CCT microstructures of L steel deformed below T'. are shown in
Figures 6.3.13 and 6.3.14. In al1 cases, the constituents were aiigned in the direction
perpendicular to the compression axis, and the y grain boundaries were not visible due to
intragranular nucleation of ae.
In cornparison with L steel as-reheated, 850 OC-deformed and 30 OC/s-cooled L
steel contained a~ ( r n d y Bm) laths which were longer and thimer and highly aiigned in
the elongated y grains. Nucleation of ae on deformation bands was also seen (Figure
6.3.13a arrow). At 10 OC/s, the degree of microstmcture alignment was slightly
decreased, and the microstructure was a mixture of asL, a B A and a ~ . The growth of a~
(mainly 83 bundles was codined to many cellular regions, as shown in Figure 6.3.13 b. a~
laths in such cellular regions were refined significantly. At 5 "Ch, the microstructure was
completely as*' with evenly distnbuted a M / y ~ . ln some areas, y grain boundaries could be
weakly seen (Figure 6.3.13~). At 1 OC/s, the microstmcture was a with some a e A and P
(Figure 6.3.13d).
A higher degree of pancahg was observed in L steel deformed at 780 OC than at
850 OC. At 30 " U s , 8-free c t ~ (B~/BI~) was the predominate constituent in an anf matrix
(Figure 6.3.14a). Compared with ae obtained at 30 " U s as described previously, a~ in
this case was wider but shorter, suggesting increased intragranular nucleation and sidewise
growth. The same changes were dso obsenred at 10 OCls (Figure 6.3.14b). At 5 OC/s the
structure was even wider B: and B? defined by the evenly distributed aM/yR (Figure
6.3.14~). At 1 "Us, the rnicrostmcture was a and a e A (Figure 6.3.14d).
The SEM CCT microstnictures of L steel deformed by double 25%~ are shown in
Figure 6.3.15. At 30 'Ch, some coarse a~ (B&:) fonned in an a~ matrix (Figure
6.3.15a). At 5 'Us, the microstructure was rnainly a B A and some a~ (Figure 6.3.1 Sb).
At 1 OC/s, the microstructure was mainly a with some a s A and P (Figure 6.3.1 Sc).
The SEM IT microstnictures of L steel deformed at 780 OC are shown in Figure
6.3.16 (fn = 1800 seconds). At 600 OC, the microstructure consisted of aeA (Br?, a and
some P, which completely obscured the y grain boundaiies (Figure 6.3.16a). At 500 OC , a
highly aiigned structure was present. 0-free a~ nucleated dong y grain boundaries or
deformation bands (Figure 6.3.16b).
6.3.3.2 M SteeI
The SEM CCT rnicrostnictures of M steel deformed below Tm are shown in
Figures 6.3.17 and 6.3.18.
In cornparison with M steel as-reheated, M steel deformed at 950 OC and cooled at
10 OCIs contained more ae that nucleated at y grain boundaries, twin boundaries (Figures.
6.3.17al) and deformation bauds (Figures. 6.3.17a2). At 5 *Ch, a mixture of GBa, P and
ae was present. The deformation bands were clearly delineated by P in Figure 6.3.12b 1.
Some a~ laths formed intragranularly to divide the prior y grain into smailer grains,
thereby refining the whole microstmcture in Figure 6.3.17b2. At 1 OC/s, only a very srnail
quantity of aeL was present.
In the case of M steel deformed at 850 OC and cooled at 10 'Ch, the
intragranularly nucleated a was dominant, and the volume fraction of a~ was low (Figure
6.3.1 8a). At 5 OC/s, P replaced the GB-nucleated aeL at y grain boundaries, and some thin
ae laths formed intragranularly (Figure 6.3.18b). At 1 OCk, the microstructure was a + P.
The SEM CCT microstructures of M steel deformed by double 2 5 % ~ are shown in
Figure 6.3.19. At 10 "Us, highly aligned aeL fonned mainly at y grain boundaries as well
as some intragranular sites (Figure 6.3.19a). At 1 OC/s, the volume fraction of aa (mainly
BE) was almost unchanged cornpared with ae in M steel as-reheated.
SEM intermpted IT microstructures of M steel deformed at 850 OC are shown in
Figure 63-20 (tn = 5 seconds). At 500 OC, the a~ transformation was cornpleted in 5
seconds, and the microstructure was aeA (Figure 6.3.20a). At 400 OC, a~ laths containing
some intralath 0 were confined to many small cellular regions (Figure 6.3.20b2).
Figure 6.3.13 SEM CCT microstructures - L steel deformed at 850 O C and cooled at: a) 30, b) 10, c) 5, and d) 1 'Ch .
Figure 6.3.14 SEM CCT microst.chires
dl
- L steel deformed at 780 OC and cooled at: a) 30, b) 10, c) 5, and c ) 1 O C k .
Figure 6. SEM CC L steel d and 780 a) 30, b)
T microstn .eforrned at OC and coa 5, and c) 1
lcture 1 O00
iled at
Figure 6.3.16 SEM IT microstructures - L steel deformed at 780 O C and held for 1800 seconds at: a) 600, b) 500 OC.
Figure 6.3.17 SEM CCT microstructures - M steel deformed at 950 O C and cooled at: a) 10 and b) 5 OCIs.
Figure 6.3.18 SEM CCT microstruc a) 10 and b) 5 OC/s.
b)
m e s - M steel defonned at 850 OC and cooled at
Figure .3.19 SEM CCT rnicrostnictures - M steel defonned at 1080 and 850 OC and coofed at: a) 10 and b) 1 "C/S.
Figure 6.3.20 SEM intemipted IT microstructures - M steel as-reheated and heId for 5 seconds at: a) 500 and b) 400 OC.
6.4 MICROSTRUCTURES OBSERVED BY TRANSMISSION ELECTRON
MICROSCOPY
Using TEM, details of the substructures and crystallographic orientation
relationships of a~ microstructures were determined. AU of the TEM observations
described in this Section were from the CCT samples.
6.4.1 Morphologies o f Bainite
as morphoiogy in L steel 1000 OC-deformed and 10 "Ch-cooled was rnainly Bm
(Figure 6.4.1 ). At the bottom of the micrograph, a slightly bending Bm bundle grew fiorn
the prior y grain boundary towards the upper left. Each aa lath consisted of "sublaths" as
those descnbed by Bhadeshia [1992]. The as lath boundaries in this bundle were clear,
but those of the bundle located in the upper nght of the micrograph were fuzzy. Note that
another a~ bundle grew fiom the pre-existing ae laths.
as morphologies in L Steel 850 OC-deformed and 10 OC/s-cooled were mainiy BI
and some Bm (Figure 6.4.2). There were many distinct cellular regions, each containing an
a~ bundle. Shown in Figures 6.4.3, ae morphologies in L steel 780 OC-deformed and 30
'Ch-cooled were aligned B: and Bi nucleated at the distorted y grain boundaries. Several
large precipitates were enclosed in the B~~ grains. Shown in Figures 6.4.4, ae
microstructure in L steel 780 OC-deformed and 10 OC/s-cooled was highly fuzzy (highly
recovered).
Figure 6.4.1 L steel defonned at 1000 O C and cooled at 10 OCIs - Overall morphology.
Figure 6.4.2 L steel deformed at 850 OC and cooled at 10 OC/s - Overail morphology.
Figure 6.4.3 L steel deformed at 780 OC and cooled at 30 OC/s - Overail morphology.
Figure 6.4.4 L steel defomed at 780 OC and cooled at 10 OC/s - Overali morphology.
ag morphologies in both recrystallized and unrecrystallized M steel were BI1 and
BIV. Two variants of intragranularly nucleated Brv laths were observed in recrystallized
M steel (1025 OC-deformed and 5 "Us-cooled) (Figure 6.4.5). n i e Brv laths were sharp-
tipped, one of which stopped growing in front of another lath without impingernent
(arrow). In Figure 6.4.6, well-developed and refined a0 bundles were present in M steel
defomed at 850 O C and cooled at 5 O C / s that contained more ae laths in a bundle
compared with that in M steel as-reheated.
Figure 6.4.5 M steel deformed at 1025 OC and cooled at 5 "C/s - Overd morphology.
Figure 6.4.6 M steel deformed at 850 O C and cooled at 5 "Ch - Overall morphology
The order-B:, BI^, Bb Bn, Bm, Brv-represents a decrease in as formation
temperature, so the TEM microstructures descried below are considered in this order.
6.4.1.1 B?
Figures 6.4.7a and b show BI^ obtained in L steel as-reheated and 1 "Us-cooled. A hi&
density of dislocations was present in the granular as. Some large growth ledges 0.03 prn
in height and 0.05 pm in spacing were observed on the BIG interface (arrow). Figures
6.4.7~ and d show an a&R island surrounded by the granular B: grains. Figure 6.4.7e
illustrates a rather fine BI^ [nicrostructure in L steel as-reheated and 10 OC/s-cooledy the
black phase was a ~ .
Figure 6.4.7 L steel as-reheated - B~'? a) B ~ ~ - i 'Ch, b) 131G - 1 'Ch (dark field), C) ~M/ .{R island - 1 d) aM/yR island - 1 'C/S
(dark field), e) BP- 10 OC/S.
Figure 6.4.8 L steel as-reheated - B/? a) B? - 1 "C/s, b) B: - 1 OC/s (dark field), c) the corresponding S M P
and indexing of a), d) B: with ledges.
6.4.1.2 B ! ~
Figures 6.4.8a and b show the details of B~~ obtained in L steel as-reheated and 1
1 laths. The OC/s-cooled. The dislocation density in B: subgrains was similar to that in B
two adjacent El: laths, al and a*, were slightly rnisonented by - 2" as shown by the
selected area diffraction pattern (SADP) in Figure 6 . 4 . 3 ~ . In Figure 6.4.8d one broad
face of B: was relatively smooth, while the other side was ragged containing ledges. The
ledges were 0.05-0.15 pn in height and - 0.6 prn in spacing.
6.4.1-3 BI
Figure 6.4.9a shows BI in L steel as-reheated and 10 'Ch-cooied. A high density
of dislocations and some s m d precipitates were obsemed Ui the as laths. a~ laths were
separated by the interlath second phase, of which some was bcc a ~ , some (in this case)
was fcc y associated with a~ in the N-W (Nishiyama- Wasseman) orientation relationship
(SADP in Figure 6.4.9b).
which is about 5-26' f?om the K-S (Kurdjumov-S
198 11.
achs) relationship porter and Easterling,
Figure 6.4.9 L steel as-reheated - BI. a) BI - 10 "Ch, b) the corresponding SADP and indexhg.
6.4,1.4_Bn
Figure 6.4.10 shows Bn obtained in M Steel as-reheated and 1 "C/s cooled. the
upper portion is the intragranularly formed Bn Coarse arc laths CO-existed with finer a~
laths or degenerate P. Bn was wavy and relatively randomly distributed with repeatedly
precipitated coarse interlath 8. The lower portion was degenerate P, slightiy etched due to
a large amount of 8.
6.4.1.5 Bq
Figures 6.4.1 la and b show Bm with intralath platelet 9 digned at - 60° to the
longitudinal axis of a~ laths. The a~ laths contained a high density of dislocations. The
SADP shows that the Bagqatski orientation reiationship Pagaryatski, I X O ] existed
between the a~ Iaths and the intralath 0 (Figure 6.4.1 1 c)
In Figure 6.4.1 Id, a high density of dislocations was present between two 8 platelets.
6.4.1.6 Bw
Lenticular Bw is unique to M steel. Figure 6.4.12 shows the details of an
intragranularly nucleated Bw lath found in M steel as-reheated. A midrib was identified
which consisted of two parallel a~ laths with an interlath phase. A group of secondary a0
laths formed at the broad face of the midnb grew in a variant - 60" to the midrib mis. As
indicated by the arrows, the right side of the Bn, lath was serrated.
Figure 6.4.13a shows an aw/ae bundle that grew fkom a GBa grain in M steel as-
reheated. The place where a~ Iaths nucleated had a hi& dislocation density (Figure
6.4.13b). This GB-nucleated aw/aB bundle evolved into a~ structure by 8 precipitation
(Figure 6.4.13 c).
Figure 6.4.10 M steel as-reheated - Bu.
Figure 6.4.1 1 L steel as-reheated - Bm. a) Bm - 10 *Ch, b) Bm - 10°C/s (dark field), c) The corresponding SADP and indexing of a), d) 8 platelets and dislocations.
Figure 6.4.12 M steel as-reheated - Brv.
In Figure 6.4.14, a fork-like aw Iath formed first at a twin boundary. Some more
a w laths nucleated fiom the broad face of the fork-like aw, foilowed by the formation of
interlath degenerate P. From the left side of the a w broad face, a BI bundle formed. As it
is well established that a w keeps a K-S relationship with y [Bhadeshia, 19921, a~ is
believed to keep the K-S relationship with y via aw.
6.4.1.8 a~
Figure 6.4.1 Sa shows a lath of auto-tempered a~ with two variants of 8 in L steel
as-reheated. Therefore a~ c m be readily distinguished from a ~ .
Unlike the auto-tempered a~ in L steel, a~ laths in M steel were free of 0 and
contain a higher density of dislocations than ae. Similar to ae growth, a~ growth was
blocked by the twin boundary (mow in Figure 6.4.1 Sb).
6.4.2 Intragranular Nucleation Sites of Bainite
In addition to y grain boundaries, the following intragranular nucleation sites of as
were identified.
6.4.2.1 Twin Boundaries
Figure 6.4.16a shows an ae bundle which formed and grew to one side of an
anneding twin boundary in M steel as-reheated. On the other side of the twin boundary,
growth of another as was blocked.
Figure 6.4.16b shows the a~ laths formed at a distorted twin boundary or
deformation band boundary in M steel deformed below Tm The kinks on the twin
boundary are noted, which are believed to have contributed to the increase of the
nucleation rate of as.
Figure 6.4.13 M steel as-reheated - GBa- nudeated aw/as. a) GBa-nucleated aw/-, b) GBa-nucleated aw/ae
(dark field), c) a~ at the growth end.
Figure 6.4.14 M steel deformed at 1025 OC and cooled at 5 "Cls - Twin boundary-nucleated awl*. a) Twin boundary-nucleated aw, b) awnucleated as.
Figure 6.4.15 a ~ . a) L steel as reheated - 30 " U s - Auto-tempered au laths with two variants of 0,
b) M steel as-reheated - 30 *Ch - a~ laths.
6.4.2.2 Preci~itates
Figure 6.4.17a shows an ae lath nucleated from a large Nb-nch precipitate in M
steel as-reheated. Figure 6.4.171, shows the interaction between dislocations and the
Nb(C,N) cluster. A dislocation pile-up was observed.
6.4.2.3 Pre-Existing- Laths
Figure 64-18 shows the nucleation of new a~ laths nom pre-existing ae laths in M
steel deformed at 850 OC and cooled at 5 OC/s. Three aspects were noted:
1 ) The nucleation of new as laths was directly fiom the broad faces of the pre-
existing c t ~ laths;
2) Mthough the new ae laths formed at different pre-existing a~ laths, they tended
to grow in the same direction (the large arrow); and
3) The at, lath 1 and 2 (the small arrows) illustrated explicitly the primary growth
stage of BIv laths.
Figure 6.4.16 Nucleation of ae at twin boundary. a) M steel as-reheated, b) M steel deformed at 850 OC.
Figure 6.4.17 Nucleation of aB at precipitate. a) M steel as-reheated - Nb(C,N) precipitate, b) M steel as-reheated - Nb(C,N) cluster and
dislocation pile-up.
Figure 6.4.18 Nucleation of as at pre-exisbg aa laths - M steel defomed at 850 OC and cooled at 5 "C/s. a) lower mamcation, b) higher magnification.
6.4.2.4 Subgrain Boundaries
Figure 6.4.19 shows two Br bundles nucleated from a curved subgrain boundary in
L steel deformed below TM. The subgrain boundary was about 0.08 pm thick. The Bi
Iath boundaries were very fùzq, representing significant static recovery.
6.4.3 Deformation-lnduced Substructures
In L steel deformed at 1000 OC and cooled at 10 OC/s shown in Figure 6.4.20, it
seemed that the dislocation density in Bm laths was higher than that in L steel as-reheated
and cooled at 10 " U s shown in Figure 6.4.1 1d. Note that the intralath ellipsoidal 0 was
coarsened and the lath boundaries were not distinct, which indicates subgrain coalescence.
Similarly, in L steel defonned at 780 OC and cooled at 10 " U s , a B: lath was
found to contain such a high density of dislocations that the lath boundary was only visible
under TEM dark field (Figure 6.4.21). Those dislocations were believed to be inhented
£tom L steel defonned below Tm by the B: lath, and the density was obviousiy higher
than that in L steel as-reheated (Figure 6.4.9a) and deformed above Tm (Figure 6.4.1).
In L steel deformed at 780 OC and cooled at 1 "Ch, a dislocation density much
lower than that in 1 O "Cls was found (Figure 64-22).
Figure 6.4.19 Nucleation of a* at subgrain boundary - L steel deformed at 850 OC and cooled at 10 "C/s.
Figure 6.4.20 Bm with fuzzy lath boundaries and eliipsoidd intraiath 8 - L steel deformed at 1 O00 O C and cooled at 10 "C/s.
Figure 6.4.2 1 Tangled dislocations in BI^ - L steel deformed at 780 OC and cooled at 10 OC/s. a) bright field, b) dark field.
Figure 6.4.22 Low density of didocations in BI^ - L steel deformed at 780 O C and cooled at 1 OC/s.
DISCUSSION
7.1 EVOLUTION OF AUSTENITE CONDITIONS WITH TMP TREATMENTS
In this section, the various types of precipitates, the kinetics of strain-induced
precipitation, the recrystallization behaviour of y deformed above TM, and the possible
substmctures formed in y deformed below Tm are discussed.
7.1.1 Precipitates in Austenite
7.1.1.1 Twes of Undissolved Precipitates
Using the solubility products in Table 2.1, the solution temperatures of possible
precipitates, Tm,., can be estimated for both L and M steels. The results are given in Table
7.1.
Table 7.1 Calculated Solution Temperatures (&J for Various Precipitates, O C
At 1 180 O C reheat temperature, it cm be seen that TiN and all Nb-rich particles are
r-
stable, consistent with the observations that both L and M steels as-reheated contain
Precipitate
0 )
?#cdt, 0
undissolved O\lb,Ti)(C,N). Although most precipitates were identified as compounds
* estimated by assuming Ti = 0.00 1 wt-%
Nb(CN) 1329 1513
NbC
1185 1318
NbN 1172 1190
VN --
1080
VC --
922
TiN 1637 1465.
Tic 953 651.
BN 1151 --
AIN
1085 924
containing two or more microalloying elements rather than simple binary precipitates,
(Nb,Ti)(C,N) cm still be regarded as a mixture of TiN and Nb(C,N) due to the different
diffisivities and solubilities of Ti and Nb in y.
7.1.1.2 Tvpes of Strain-Induced Precipitates
Ti has the strongest af5nity to N among al1 elements in the steels investigated by
the stochiometric ratio of Ti : N = 48 : 14 (atomic weight ). It is believed that TiN foms
first at the reheat temperature, and subsequently M(C,N) or V(C,N) foms during
deformation or cooling.
ln L steel, the N available to form Nb(C,N), N*, is N - 14TV48 = 0.006085 (wt-
%). Therefore, the formation of Nb(C,N) is possible.
[n M steel, the stochiometric ratio of Nb : N in atomic weight is 93 : 14 = 6.64,
higher than the actual weight percent ratio of Nb : N = 0.065 : 0.0147 = 4.42. Although
Ti was not detected due to the limit o f the chemicd composition analysis, it was often
present in precipitates, so Ti content was taken as 0.001 wt-%. Hence N' = N - 141248
= 0.0147 (wt-%), almost equal to the initial N content. Obviousiy Nb(C,N) could be
formed.
For sirnplicity, ordy the precipitation behaviour of the equilibrium state of Nb(C,N)
were estimated after the method provided by Speer et al. [ 19871 (Appendix). The results
are given in Table 7.2. Here f is the precipitate volume fraction, AG and AGp are the
chernical driving force for nucleation of Nb(C,N) and the total fiee energy change for
precipitation, respectively, as shown in Equations 2.6 and 2.7, respectively. ml, [Cl and
N are the concentration of each element in solid solution, and k, is the saturation ratio.
Table 7.2 Calculated Precipitation Parameters of Nb(C,N) at 1 180 OC for L and M Steels
It is noted that: 1) for L steel as-reheated, N = 2.083 x IO'" means that essentiaüy
al1 N is combined with Ti and Nb as undissolved (Nb,Ti)(C,N) particles. Thus, the strain-
induced precipitates should be carbides (i.e (Nb,Ti)C d e r than carbonitrides and 2) the
Terms
f(%) [Nb] (w-w [Cl (wt-%) pl'] (wt-%)
k AG (J/md) AGp (J/mol)
higher k, and AGp of M steel means a stronger tendancy of precipitation than of L steel, which
explains the higher density of strain-induccd precipitates associated with M steel. In M steel,
the strain-induced precipitates were mainiy (Nb,Ti,V)(C,N) or (Nb,V)(C,N) and some
Nb(C,N) and V(C,N). No AIN was observed in either steel. This may be attributed to its
high solubility product in y in the temperature range 900- 1 500 O C [Sumki et al., 19831.
* Calculated on the assumption that Ti takes up N fkst.
L Steel 4.847 x lo4 0.05 17 O. 1048
2.083 x 10-l4 2.678
- 3.23 x lo4 - 15.66
M Steel 6.716 x lo4
0.009 1 0.433 1 0.006273
7.184 - 8.68 x 10' - 58.28
7.1.1.3 Mechanisms of Strain-Induced Precipitation
There are four types of sites for the nucleation of strain-induced precipitates: 1)
prior y grain boundaries and subboundaries, 2) defonation defects (dislocations and
vacancies), 3 ) twins and deformation bands, and 4) pre-existing precipitates.
Defonned M steel contained a high density of strain-induced precipitates: the
lined-up precipitates were medium-sized (- 20 nm), the scattered precipitates were fine (-
5 nm) (Figure 6.1.4). Dimensions of the areas outlined by lined-up precipitates are of the
order of microns, suggesting that strain-induced precipitates are distributed at subgrain
boundaries or prior y grain boundaries. The faster diffusion rate of solute atoms dong
these boundaries results in larger precipitates cornpared with precipitates scattering in the
matrix.
Also, the strain-induced precipitates tend to be present in clusters as show in
Figures 6.1.3 and 6.1.6. It is believed that the clustenng of precipitates is an actual
phenornenon. Zou and Kirkaldy [1989] attnbuted it to the over-etching of the matrix, and
Nelson [1996] thought it just a replicating effect. However, the frequent observations of
precipitate clusters in TEM replicas lead to the belief that the clustering of strain-induced
precipitates is an important mechanism in this study. Using the Crz03-based electrolyte
polishing solution when preparing thin foils, the precipitates were well preserved and
precipitate clusters were observed in thin foils. For example, Figure 6.4.17b showed a
precipitate cluster in M steel. The mt-chanisms of precipitate clustering is probably due to
the physical presence of coarse particles such as MnS in L steel and (Nb,Ti)(C,N) in M
steel, because coarse particles usudly have surface defects that can assist the nucleation of
new precipitates. A Nb depletion area is expected around the coarse (Nb,Ti)(C,N) during
its growth, which is unfavourable for nucleating new Nb-rich precipitates. On the other
hand, deformation builds up a high density of dislocations around these coarse partictes as
observed by TEM, which facilitates the diffusion of solute atoms such as Nb and Ti
towards the pre-existing coarse particles through dislocation pipe diffision, favouring
precipitate clustering.
There are also fine precipitates that could form at dislocations or vacancies in the
matrix especially in M steel (Figure 6.1.4). The vacancies are produced by either
deformation or non-equilibrium segregation [Jonas 19881 (e.g., Nb and B) during cooling.
7.1.1.4 Kinetics of Strain-induced Preci~itation
For Nb-bearing steels, the time for 5% strain-induced matnx precipitation, hoj, is
calculated using Equation 2.8 developed by Dutta and Sellars [1987]. Here the constants
A and B are determined as 3 x lo6 and 2.5 x Io1', respectively, according to Dutta and
Sellars [1987]. Figure 7.1.1 shows t0.05 as a fbnction of Td for L and M steel, respectively.
For L steel defomed at 1000 OC, 5% precipitation appears in less than I second,
while at 850 or 780 OC, 6-22 seconds are required. This explains the fact that there are
many strain-induced precipitates in L steel deformed at 1000 O C while very few such
precipitates in L steel deformed at 850 or 780 OC.
For M steel deformed above 1025 OC, 5% precipitation occurs in less than 10
seconds, while at 950 or 850 OC deformation, hos becomes 35 or 100 seconds. Therefore
the caiculation is consistent with the experimental observations that there was a very high
density of strain-induced precipitates in M steel deformed at 1080 OC but a lower density
of precipitates in M steel deformed at 950 OC, and even lower at 850 OC.
Figure 7.1.1 5% precipitation time ( h ~ ) W. defonnation temperature ( la: a) L Steel, b) M steel.
7.1.2 Recrystallization of Gustenite
Dynamic recrystallization occurs when the applied strain (E) exceeds a cntical
a
value E,.. By Equation 2.10, E, of L steel is 0.85 for a main rate E of 1 s-' and Td of 1000
a
O C (highest); E, of M steel is 0.89 for E of 1 s" and Td of 1 080 OC (highest). The applied
50% strain E = ln ( L A ) = 0.69 < E, for both L and M steels. Here L is the original
sample length and Ld is the deformed sample length. Therefore for deformation above TNR,
static recrystallization rather than dynamic recrystallization could occur.
The time for 50% static recrystdlization, las, for various Td were calculated
according to Equation 2.12. The results are given in Table 7.3.
Table 7.3 Calculated 50% Recrystdization The, t0.5, of Austenite Deformed Above T V R
Td, OC 1180 1080 1025 1 O00 950 850 780
to.5 sel L Steel 0.0002 0.00 17 0.0060 0.0 1 O6 0.23 88 20.03 50.54
M Steel 0.00 1 1 O .O079 O. 0270 0.0488 1.100 92.00 232.1
Hence, both deformed L and M steels can recrystallize rather rapidly when Td >
950 O C , but rather slowly when Td c 850 O C . SO the experimentally measured TM* of 950
O C for L steel matches well with this calculation. For M steel, the rneasured Tm of 1025
O C is 75 O C higher than the calculated 950 OC. The discrepancy may be due to the
significant grain-boundq pinning effect applied by the fine strain-induced precipitates
distributed at y grain boundaries, which raises the TM of the M steel. However, if Nb
remains undissolved because of a low heating temperature, it does not demonstrate any
delaying effect [Cordea and Hook, 19761. To estimate TNR using Equation 2.9, the
amouat of Nb in Nb-nch undissolved particles should be subtracted £iom the overall Nb
content.
Re-applying the values of w] and [Cl in Table 7.2 to Equation 2.9, new Tm
temperatures for L and M steels are 1047 and 1002 O C , for L and M steel, respectively,
matching well with those found experirnentally (950 and 1025 OC).
7.1.3 Substructure of Austenite
In this study, the y substructure is determined indirectly by observing the
transfo nned microstructures.
As illustrated in Figures 6.3.13, 63-20, and 6.4.2, the boundaries of the cellular
substructures are revealed by the boundary-nucleated ae laths. The boundary thickness of
these cellular substructures is - 0.08 Vrn (Figure 6.4.1 5 b), suggesting that the boundaries
consist of a high density of dislocations.
0
Roberts et al. [1978] developed an equation, aven a constant strain rate E = Us,
to estimate the subgrain size, Dd (m), in low carbon steels which is dependent on the
deformation temperature Td (K)
0, = 269 exp (-4770 / Td)
The calculated values of Dd are 3.8 and 2.9 pm, respectively for 850 and 780 O C -
deformed steels. By direct measurements to L and M steels 850 OC-deformed, the average
diameters of the cellular regions are 3.6 and 3.2 pm, respectively. Some recent
publications consider the intragranular nucleation of a~ to be mainly deformation bands
[Yamamoto et al., 1 9951 or dislocations mjiwara el al. 19951. However, on the basis of
the fact that the cellular substnxcture diameter is at the same magnitude as the calculated
subgrain size, and the smoothly curved boundaries are different from the çtraight twin
boundaries or deformation bands, it can be concluded that these cellular substnictures are
dislocation subgrains. Each subgrain acts as an isolated grain [Araki, et al., 197 1 ; Baiiey,
19631.
Considenng the standard 10 s holding after deformation and B., the formation of subgrains
starts at - 15 s, and completes at - 20 s in L steel defomed at 850 OC (Figure 6.3.13).
For L steel deformed at 780 OC, the subgrain boundaries are difficult to see because of a
prevailing intragranular nucleation. For M steel, substructures are present at - 400 O C IT
for M steel defomed at 850 OC when GBa is avoided.
7.2 VARLATION IN BAINITE CCT KINETICS WITE AUSTEMTE CONDITION
The CCT resuits in Section 6.2 show that a~ transformations are slightly
decelerated due to deformation above TNR but significantiy accelerated due to deformation
below TNR.
7.2.1 Bainitestart Temperature
For IT, B, is the highest holding temperature (IT BJ at which arc starts to form at
a detectable rate and above which no ae c m be observed. For CCT, B, is usually defined
as the highest temperature (CCT B.) at which a deviatioa on the L-T curves is detected by
the LVDT of the dilatometer when about 5% a~ has formed. Therefore a CCT B, is
determined by the rate of nucleation and growth of ae, and is always lower than an IT 8,.
The CCT Bs can be directly rneasured by the dilatometer as shown in Section 6.2. The IT
Bs can be estimated by examining the SEM microstmctures in Section 6.3.
For L steel as-reheated, the IT Bs is found to be between 600-700 O C , and the
CCT B, (aeA) is 6 1 5 O C . Considering the calculated IT B, of 660 O C using Equation 3.1
[Steven and Haynes, 19561, the tT B. (aBA) is taken to be 660 OC. Similarly, the CCT B,
(aBL) is 510 OC and the IT B, (aeL) is above 500 OC. Assuming the IT Bs is higher than
the CCT B. by 660 - 6 15 = 45 OC, the IT B, (aeL) is: CCT Bs (aeL) + 45 = 5 10 + 45 =
555 OC.
In L steel defomed at 1000 OC and cooled at 1 OCfs, the CCT B, ( a B h ) is 580 OC,
the IT B. (aBA) above 600 O C , and the IT B. (aBL or Bm) above 500 OC. Keeping the same
difference of 45 OC between IT B. and CCT B., the IT B. (aeA) is taken to be 580 + 45 =
625 OC. At 5 OC/s, the CCT inicrostmcture is aeA (B: and B:), but the 520 OC-
intempted CCT microstmcture is mainly a~ and ~f plus very few B: (Figure 6.3.7c),
which means that the CCT B. (@) is - 520 OC, so the IT B. (B[~) is - 520 + 45 = - 565
O C .
For L steel deformed at 780 OC, the main constituents are and B: with a CCT
B. of - 600 OC at 10 OC/s, or is Br with a CCT B. of 410 OC at 30 "C/s; by IT of 1800
seconds, the IT Bs (Bit) is above 650 OC, and the IT Bs (Bi) is above 500 OC. So the
difference between the IT B. and the CCT B, is more than 50 OC.
For M steel as-reheated, the IT B, is found to be between 500 and 600 O C , and the
CCT B, is 520 O C by examining the SEM microstmctures. The calculated IT B, is 570 O C ,
therefore the IT B, is taken to be 570 OC.
For M steel deformed at 1080 OC, the measured CCT B, is 425 O C but a smali
amount of as emerges in 520 OC-intempted CCT microstructure. Therefore 520 O C is
regarded as the IT Bs.
For M steeI defomed at 850 OC, the IT B, is above 600 OC. However, the CCT Bs
is 520 OC, which means that the difference between the CCT B. and the IT B. is more than
70 OC.
Table 7.1 summarizes the results of IT Bs and CCT 8, for both L and M steeIs.
Table 7.1 Bainite Start Temperatures, OC
L Steel Bs As-Reheated Td = 1000 O C Td = 780 OC
Ir Bs @aA) 660 625 > 650 TT Bs ((laL) 555 - 565 > 500
CCT BXU& 615 580 600 CCT ~ . ( a e ~ ) 510 520 410
M Steel Bs As-Reheated Td = 1080 O C Td = 850 O C
IT B, 570 520 > 600 CCT Bs - 520 525 520
For L steel, the IT B,'S (ae? are consistent with the CCT ElsWs. Also, deformation
has little effect on B, of aaL.
For M steel, the IT B,'S have the same trend as for L steel although the CCT B.'S
keep relatively constant. The reason is due to the presence of GBa that compensates the
increase in IT B, for an increased overall C concentration.
7.2.2 Bainite Continuous Cooling Transformations
L steel as-reheated has a y grain diarneter of 21 pm and a very low density of
undissolved (Nb,Ti)(C,N) precipitates. Most Nb and al! B are in solution. The theory that
B and Nb segregate to y grain boundaries dunng cooling is indirectly evidenced by the
observation that a and P transfomations were postponed to a longer time dut-hg
continuous cooling (Figure 6.2.3a), and a* was obtained over a wide range of cooling
rate.
At fast cooling rates 130 - 30 "Ch, a large driving force is provided by large
undercooling. aa can fonn at y grain boundaries and proceeds very rapidly so that
expansion in dilatometric sample length (AL) is very large within a very small temperature
range even though the interfaciai energy of y grain boundaries may be reduced by B-Nb
segregation (Figure 6.2.4a).
Slow cooling rates, 5-1 "Ch, ailow sufficient time for the de-segregation of non-
equilibrium B-Nb fiom y grain boundaries [He et al., 19881. so a~ may nucleate at a
srnaller undercooling (i.e. at higher temperatures), resulting in an increased CCT B, as
shown in Figure 6 .2 .4~ . A raised CCT B. usually lads to an increased ae volume fraction
when no a or P fonns before a ~ .
7.2.3 Deceleration of Bainite Continuous Cooling Transformation
7.2.3.1 Recystallized L Steel
While a transformation is accelerated, a~ transformation is deceterated over 30-1
"C/s due to four factors influencing both nucleation and growth rate of as (Figure 6.23b):
1 ) Recrystallized y provides a slightly refined y grain (fiom 2 1 to 17 pm) or more
nucleation sites for all new phases because the "effective" y grain surfaces for nucleation is
assumed to be proportional to the total y grain boundary area phadeshia, 19921.
Aithough codicting results are reported on the effects of decreased y size on a3
transformation kinetics pmemoto, et al., 1980; Yamamoto, et al., 19951, the effects of
the y grain size refinement on the nucleation of a~ in this study should be enhancing as y
grain boundaries are the main nucleation sites of a ~ . But the magnitude of this enhancing
effect is small since the net increase in y surface area is limited.
2) B resegregates to the "fiesh recrystdlized y grain boundaries rapidly, so GBa
is not present in L steel defomed above Tm. A s the resegregated B concentration is
decreased due to increased total y grain boundq are& ae nucleation is considered to be
favoured, but this effect, if any, is fairly srna11 because the refinement of y grains is small.
3) The refinement of y grain size due to recrystdlization shortens the final length of
a0 laths.
4) In L steel defomed above Tm, the presence of many strain-induced (Nb,Ti)C
precipitates aiters the chemistry of the pnor y, and thus affects a0 growth.
In L steel, as ae nuclei occupy al1 the y grain surfaces before they grow to 5%
volume, B. is determined only by the growth rate of those nuclei as a result of nucleation
site saturation. In this case, Morozov and Vovokhov [1989] proposed a simple equation
to describe a~ incubation time t as:
where X is the volume fiaction of ae, D the y grain size, and G the ae growth rate. It is
clear that Bs (i.e. the value off when X is about 0.5%) is proportional to the y grain size
and the inverse of the growth rate of ae.
The isothermal growth rate G is estimated frorn a formula developed by Yoshie et
ai. [1988]:
[Nb])G, G = k, [Nb] e ~ p ( ~
where k-r and ks are coefficients representing the effect of w] on the suppression of a~
growth, Gk is the growth rate of bainitic a in carbon steel [Kauhan, 19621.
It can be seen that G decreases exponentially with the decrease in w] due to
strain-induced precipitation, leading to a decrease in B.. This effect must be so large that
even the increase in y grain surfaces and the decrease in resegregated B concentration
cannot balance it.
At 5 to 1 OC/s where the non-equilibrium segregation of B is able to disappear, the
increase of y grain boundaries leads to more a and fewer aeA. Because cooling rate and B
segregation are related to the driving force for ae transformation, it seems that a~ foms
by a displacive transformation.
7.2.3.2 Unrecrystailized L Steel
As softening is influenced by the chemistry of the steel, particularly the
rnicroalloying elements such as Nb, the initial grain size and the preceding strain, the
sofkening ratio should be very low in this case under the low deformation temperature, an
unchanged @%] and fast cooling. Therefore at fast cooling rates 130-30 "Us, the work-
hardened L steel is mechanicaily stabilized, causing a decrease in BJM, by 25-55 OC.
Deformation at 780 OC causes a larger decrease in BJMs than deformation at 850 O C due
to a higher wok-hardening at 780 OC.
7.2.3 -3 Recrystallized M Steel
In M steel as-reheated, GBa always covers the pnor y grain boundarïes within the a~
region. Although some aw/ae fonns at GBa, ae mainly nucleates intragranularly,
therefore y grain boundaries are not as important as those in L steel. In this case where
there is no site saturation, an equation that integrates the dimension, nucleation rate and
growth rate developed by Morozov and Vovokhov [I989] is applicable:
where L is the final dimension of aB, which is always smaller than the average y grain size
for intragranuiar nucleation and growth, I the nucleation rate of a~ nucleus per unit
volume of y, m the ratio of half-width (o) to length of a~ lath (L).
Cornpareci with M steel as-reheated, M steel deformed above TNR has the a~ laths
that are shorter in length but constant in width. Therefore Equation 7.4 can be rewritten
as:
Hence, r for 5% a~ or CCT B. is detemined by the nucleation rate I and the
growth rate G, which are afTected by two factors:
1) With the introduction of more GBa due to deformation above Tm, the C
concentration in the remaining y is increased, so the a~ phase is hard to form and I is
decreased.
2) Simiiar to that in L steel deformed above TM, the decrease in [Nb] because of
the strain-induced Nb-rich precipitation decreases G.
Note that the strain-induced precipitates tend to form at y grain boundaries and
undissolved coarse particles, which promotes, however, the nucleation of both a~ and a.
Therefore B. is decreased with a corresponding decrease in a* volume fraction.
7.2.4 Acceleration of Bainite Continuous Cooling Transformation
7.2.4.1 Unrecry st allized L Steel
In the intermediate cooling rate range of 10-5 "C/s, a~ transformation is
significantly accelerated due to:
1 ) increased nucleation sites such as - 25% increase in y grain boundary areas by
their elongation, the incoherent annealing twin boundaries, deformation bands and
subgrain boundaries.
2) increased stored energy in terms of dislocations and vacancies especiaily in the
vicinity of the boundaries mentioned in I ).
3) diluted B concentration at these boundaries.
It is very interesting that the increase in aBA compensates for the decrease in aBL,
hence the total ae volume fraction is not obviously influenced by deformation below Tm.
7.2.4.2 Unrecrvstallized M Steel
The factors influencing a~ transformations in M steel deformed below Tm are:
1) Increased nucleation sites for a* and stored energy sirnilar to those in L steel
deformed below TNR.
2) Increased overall C concentration in the remaining y due Iargely to introduced
GBa. The enhancing effects due to increased nucleation sites and stored energy to as
transformation is not as significant as in L steel.
3) Decreased a~ growth rate due to strain-induced precipitation.
So the combination of the above three factors only causes a slight acceleration of
a~ transformation in M steel.
7.3 BAINITE NUCLEATION SITES
Through d l the observations, ae is found to nucleate heterogeneously rather than
homogeneously. In Sections 6.3 and 6.4, the heterogeneous nucleation sites for a e are
identified as: 1) grain boundaries and twin boundaries, 2) precipitates and pre-existing aa
laths, and 3) deformation substructures. These aa nucleation sites are discussed
individually in this Section.
7.3.1 Grain Boundaries and Twin Boundaries
7.3.1.1 Recrystallized Austenite
In recrystallued L and M steels, u e normally nucleates at y grain boundaries when
there is no GBa.
For the y + a transformation, the fiee energy change for heterogeneous
nucleation, AG,,,, is given by [Poner and Easterling, 198 11
AGher = - VAG, + VAGl i- AR AGd (7-6)
where VAG, and VAG, are, the reduction of the volume fiee energy AG, and the increase
of the misfit strain energy AG, respectively, due to the creation of a volume V of a. AR is
the interfacial 6ee energy R increase due to the creation of an area A. AGd is the released
free energy when the creation of a nucleus results in the destruction of a defect.
Therefore grain-boundary nucleation of ae is energetically favoured because of the
destruction of the boundary area (high angle interface with high energy [Nabarro, 19881)
covered by an a~ nucleus. Further, a~ assumes the K-S or the N-W orientation
relationship with y (Figure 6.4.9), thus making the nucleation easier.
As annealing twin boundaries are usually CO herent boundaries with low int erfacial
energy wetals Handbook, 19851, a large undercooling must be needed for nucleating a*
at annealing twin boundaries. In the case of deformation above-TNR, deformation causes
the annealing twins to Iose coherency with y [Tamura, 19881, and some twins should
remain incoherent when y is yet not fully recrystallized. According to Porter and
Easterling [1982], the interfacial energy of incoherent and/or semi-coherent anneding
twins is comparable to that of grain boundaries.
The experimental results are consistent with this analysis because nucleation of ae
at twin boundaries occurs only in fast cooled L steel. Examples are L steel as-reheated
and 130 OC/s-cooled (Figure 6.3. la), or deformed at 1000 OC and 30 OC/s-cooled (Figure
6.3.6a). For M steel, examples are low-temperature transformed samples such as
defomed at 1080 O C and 5 seconds IT at 500 O C (Figure 6.3.12b).
In contrast, in recrystallized L steel 10 OC/s-cooled (Figures 6.3. lc and 6.3 -6b) and
recrystallized M steel 1 OCls-cooled (Figure 6.3.3 c) where a complete recrystailization
occurs and a small undercooling is available, anneaiing twin boundaries are not the
preferred a~ nucleation sites. Obviously, as nucleation is dependent upon y condition and
undercooling, Le., free energy change for the transformation.
For M steel during CCT, some ae nucleates at GBa in the as range (1 0- l OC/s) or
from hÿin boundary-nucleated a (Figure 6-4-14), which is classified as aw by Ohmori
[1991]. Such a definition may be applicable when the high cooling rate (e-g., 10 'Us)
suppresses the 8 precipitation (Figure 7.3.1 a). However 5 "Us ailows the intralath and/or
interlath 8 precipitation to occur at the bundle growth ends, so such a microstmcture
should be properly described as ae rather than aw (Figure 7.3. l b). Such a microstructure
is also observed in TEM as shown in Figure 6.4.13. Therefore in tbis study, aa and aw
are regarded as the same phase during CCT.
GBa becomes thicker with the decrease of cooling rate or Td because a slower
cooling rate favours C diffusion for GBa formation, and a lower Td results in more refined
y grains. Additionally, the medium sized strain-induced precipitates distributed on both y
grain boundaries and twin boundaries encourage the nucleation of GBa since grain
boundaies with inclusions have lower energy banier for nucleation than grain boundaries
b mis on, 19821. Aithough the increase of GBa provides more nucleation sites for GBa-
nucleated a*, it decreases the overd a~ nucteation by raising the overd C concentration
of YR.
Figure 7.3.1 GBa-nucleated a ~ . M steel deformed at 1025 O C and cooled at: a) 10 and b) S°C/s.
In unrecrystallized y, the rate of a~ nucleation at y grain boundaries and twin
boundaries is enhanceci because:
1) The y grain surface area is increased by - 25% due to an ehpsoidal y grain
forrned by the 50% deformation below-Tm [Tamura, 19881, and y grain boundaries are
distorted, thereby the potential of y grain boundaries for the nucleation of ae is increased.
2) Deformation below Tm results in incoherent and distorted amealing twins as is
shown by TEM microstructures in Figure 6.4.2 la where the twin boundary contains steps.
7.3.2 Precipitates
Nucleation of ae on the coarse Nb-rich precipitates as s h o w in Figures 6.3.4 and
6.4.17 may be attributed to two factors:
I ) Around the coarse Nb-nch precipitates or precipitate ciusters, a region of Nb-
depletion is created as difisivity of Nb is much lower than that of C and N, and thus a~
nucleation is chemicaily favoured in çuch a region.
2) A stress concentration is built around coarse precipitates or clusters as a result
of deformation and "differential thermal contraction dunng cooling that could give rise to
plastic strain in the matrix and a high dislocation density" [Barritte and Edmonds. 19821,
as indirectly observed in a (Figure 6.4.17). Within the stress field, C and Nb atoms tend
to stay in the compressive side of dislocations, thus creating a C- and Nb-depletion area
that is favourable for a~ nucleation.
L steel as-reheated or deformed below Tm has ody few fine undissolved
(Nb,Ti)(C,N) particles, so they have no significant role in nucleation of u ~ . L steel
deformed above KvR at 1000 OC contains a high density of strain-induced NbC or (Nb,Ti)C
precipitates. The radially distnbuted ae laths in Figure 6.3.6a2 suggest that nucleation on
precipitates is possible.
M steel as-reheated has few coarse undissolved O\ib,Ti)(C,N) particles, and M
steel deformed above TM or below TNR has a high density of strain-induced (Nb,V)(C,N),
Nb(C,N) and V(C,N) precipitates. Both groups of precipitates are observed to nucleate
a* (Figures 6.3 -4 and 6.4.1 7).
7.3.3 Pre-Existing Bainite LathsiPlates
A theory of either autocatalytic nucleation [Olson and Cohen, 19811 or
sympathetic nucleation [Aaronson and Wells, 19561 cm account for the nucleation of a e
on pre-existing ae laths from the broad face of the pre-existing ae laths (Figure 6.4.1 8). It
should be the physical presence of pre-existing ae laths that creates a stress field and then
encourages the formation of new a* laths.
7.3.4 Deformation Substructures
Deformation substructures such as deformation bands and subgrain boundaries
contribute to the increased nucleation sites and rate. Both defonnation bands and
subgrain boundaries contain a high density of dislocations, which can assist the nucleation
of a~ in a way similar to y grain boundaries.
Generaily, substmctures are mostly eliminated in recrystallized y. For L steel
deformed at 780 OC, the subgrain boundaries are difficult to see because of a prevailing
intragranular nucleation.
Subgrain boundaxy nucleation does not occur when the cooling rate is too fast or
too slow. For example in L steel defomed at 850 "C and cooled at 30 *C/s where a* can
grow across the whole the y grain without the confinement of subgrain boundaries; or
cooled at 5 OC/s where subgrains may have disappeared due to dislocation movernent and
the microstructure is carbide-fiee aBA due to sufficient C diffusion. Only at 10 OC/s does a
well-developed subgrain structure forms.
In M steel CCT, no obvious subgrain boundaries are seen. in M steel IT, subgrain
boundary-nucleation of a~ becomes apparent at lower ( e g , 400 O C ) rather than higher
temperature (cg., 500 OC) as shown in Figure 6.3.20.
7.4 GROWTB OF BAINITE
7.4.1 Bainite Lengthening and Thickening
In most cases a~ has a lathlike morphology, indicating that the growth rate is
higher in one direction. It is considered that the lengthening of a~ is separated from the
thickening of ae because:
1) By interrupted IT, it is seen that aa lengthens to its final dimension in times of
order of seconds, but the a~ thickening continues dunng further holding (Figure 6.3.12).
2) By CCT, it is seen that the length of the a~ laths in L steel as-reheated is
relatively constant over cooling rates of 1-130 "Us, while the width of the a~ laths
decreases with increasing cooiing rate. Note the drarnatic decrease in a~ lath length
O C C U ~ ~ ~ at 1 "Cls due to the formation of a (Figure 7.4.1).
It has long been argued whether asL is saturated with C dunng ae growth. In this
study, it is assumed that ae laths are at least partially saturated with C, thus 0 precipitates
directly from a* laths. Evidence is that the intralath 8 maintains the Bagarayatsky
relationship with the ae laths as s h o w in Figures 6.4.1 1c. On the basis of this
assumption, as lengthening occurs rapidly below B, by a displacive mechanism with C
being trapped in the a~ laths. Such a lengthening cannot cross grain boundaries, twin
boundaries, and, in the cases of L and M steels deforrned below Tm, subgrain boundaries.
This mechanism is different from diffusion-controlled a growth where a can grow across
all these boundaries.
For thickening, Br cm grow into the wider B: only at slow cooling rate and higher
temperature because the latter needs more time for the difision of the substitutional and
interstitiai atoms. The superledges found on the broad face of B: strongly imply that aa
thickens through a Ledge growth mechanism (Figure 6.4.8d). At relatively lower
temperatures, ae thickens mainly by a coalescence process in which atoms migrate
through the adjacent smaü angle a~ Iath boundaries. Evidences are the observations of the
small angle boundary and coarsened ellipsoidal intralath 8 associated with Bm, the loss of
sharpness of lath boundaries (Figure 6.4.1 l ) , and the coarse a~ microstructures of long
time IT microstmctures shown in Figure 4-3-12. Further, the slight misorientation of 2"
(Figure 6.4.1.8) of two neighboring ae laths also suggests that a coalescence by slight
rotation of adjacent laths be possible. Therefore a~ thickening is a slower process by
mainly ledge mechanism at higher temperatures and mainly coaiescence at lower
temperatures.
Figure 7.4.1 The variation of bainite dimensions with cooling rate in L steel as-reheated. a) length vs. cooling rate. b) width v-r cooling rate.
7.4.2 Evolution of Bainite Morphologies
As descrïbed in Section 6.2, a~ has two major a morphologies: lathiike and
lathless. The distribution of secondary phases (O and aM/yR) dictates the secondary
morphologies.
Lathiike a* (aaL) nucleates sympatheticdy and grows displacively into bundles.
However, the formation of BIv is by the growth of individual laths according to the
following steps:
1 ) A midnb forrns first as shown in Figure 6.4.18.
2) The sidewise growth of secondary a~ laths fiom the broad face of the rnidnb
becomes possible after C diffuses out of the C - e ~ c h e d midrib. Those secondary a~ laths
are short probably due to C saturation of the surrounding y. C again diffises out of the
secondary ae laths into y. These secondary a~ laths are shown by the serrated edges in
Figure 6.4.12. The self-accommodation and the migration of boundaries of those
secondary as laths mate the outline of a Brv crystal smooth and Ienticularly shaped.
M steel deformed at 1080 OC and 1800 seconds IT at 400 O C obtains an
incomplete as microstructure even though 5 seconds IT has transformed most y into ae
(Figure 6.3.12). This suggests that the a~ transformation cannot proceed when C-
enrichment of y reaches some level.
For lathiess ae or aeA (B: or ~ ~ 9 , BIG is considered to form through a difision-
controlled mechanism because of the highest a~ formation temperatures where such a
mechanism is possible, the ledges on its broad face (Figure 6.4.7), and its ability to grow
across y grain boundaries.
During a~ growîh, there are two competing mechanism of C difision: one is the
partitioning of C corn the C - e ~ c h e d a~ phases to retained y and another is the
precipitation of interphase 0 [Bhadeshia, 19921. To form BIG, a long-range C difision is
required so that C atoms can diffise over one or more ae grains to form aM/yR. BI
requires that C partitioning outweighs 0 precipitation so that C cm diffuse to the
surroundhg y. B: means a coalescence and sidewise growth of BI. Ba forrns when the C-
enrichment in y exceeds its limit (i.e., about 2.2% [Krauss, 19881). Bm demonstrates that
the 0 precipitation dominates. Brv has either interlath and intralath 8 or a& dependent
upon the secondary growth. Considering that the formation temperature descends in an
order of B:, BI^, BI, Bn and Bm and BR, long-range C diffusion usually occurs at higher
ternperatures.
As C difisivity in a is much higher than that in y [Bhadeshia, 19921, the latter
(D7c) is the controlling factor. According to Wells and Mehl [1976], D>c (cm2/s) is given
by
-3 2,000 Dc = 0.12 exp RT
Assuming that the flux of C is one-directional dong a coordinate z normal to the
a/y interface, a mode1 to explain the partitioning of C f?om supersaturated % was
proposed by Bhadeshia [1988] as
where t d is the time required to decarburise the supersaturated ferrite, o is the asL width,
X is the average mole fraction of carbon in the dloy, 5 is the weighted average diffusivity
of C in y, and x4 and ga are the paraequilibriurn C concentration in a and y, respectively.
It is clear that td CC 115. Although O of BI is 1.5 tirnes larger than o of Ba with a
decreasing transformation temperature, t d increases exponentidy with the exponential
decrease in D, leading to a replacement of BI by Bm.
The influence of y condition on aB morphologies are then discussed below.
7.4.2.1 Recqstallized Austenite
In L steel as-reheated, a~ appears mainly as Bm at higher cooling rates (i-e., 30-5
'Ch) with Br being the secondary phase, and as aBh (B: f BI^) at slow cooling rate (i-e., 1
"Ch). Bn is rarely present because during CCT of L steel, C-enrichment in y~ is oot
saturated to the limit of interlath 0 precipitation.
No significant changes in a~ morphologies occur in L steel deformed above Tm at
1000 OC.
In M steel as-reheated, GBa-nucleated aa grows into lathlike BI or Bm bundles
depending on the cornpetition between C partitioning and 0 precipitation. Al1
intragranularly nucleated a~ is lathlike Bnr. It is noted that B, rises with the decrease of
cooling rate and hence intralath 0 decreases or disappears.
In M steel deformed above Tm, it is not surprishg that B ~ v is shorter and thimer
than in M steeI as-reheated as the decrease of y.gr"n sizes and the significant suppression
of the growth of the secondaq ae lath perhaps due to higher overdl C concentration and
thus lower transformation temperatures as a result of more GBa.
7.4.2.2 Unrecrystallized Austenite
The striking feature is that Bm is largely replaced by Bi or B: (850 OC-
deformation) or eliminated (780 OC-deformation) in L steel defomed below Tm, saying
that the C partitioning cornpetes over the û precipitation because of
1) Dislocation pipe diffusion provided by the deformation-produced dislocations.
It is further supponed by the fact that Bm is compfetely eliminated by decreasing T d h m
850 OC to 780 OC because a lower Td normaily results in more dislocations [Bhadeshia,
19921.
2) Raised Bs due to reasons listed in Section 7.1, which favours a quick C
diffision.
Note that the intralath 0 aiways keeps one variant with a* laths. In other words,
the intraiath 0 should precipitate out at the closely packed planes of y, which is confirmed
by the TEM observation where a high density of accommodation dislocations is presumed
to be introduced after the precipitation of intrdath 0 platelets (Figure 6.4.1 le) Othenvise
the intraiath 0 should precipitate preferably a t those dislocations to fom randomly
distributed variants. The pnor y grain is divîded by severai subgrains. Unlike a
recrystailized y grain that nonnally contains several as bundles that impinge each other,
each subgrain usually has only one c r ~ bundle completely confined to the subgrain. This
kind of refinement of y grain is considered similar to that of aM [Ohrnori and Mala, 199 1 ;
Tsuzaki rf aL, 19911.
For IT M steel, it is worthwhile to compare the IT microstnictures obtained in M
steel deformed above and below T N ~ (Table 7.5)
Table 7.5 Cornparison of Partiaily IT Microstructures of M SteeI Defomed Above and Below TNR
IT Treatment
At 500 OC holding, a significant acceleration of as transformation is redized in
unrecrystallized M steel. Obviously, unrecrystallized y of M steel provides more driving
force for a~ transformation due to stored energy that favours the formation of aeA rather
than a a L and promotes C partitioning. At 400 OC, ae transformation proceeds rapidly in
both steels because the undercooling is larger than the deformation introduced energy.
However, the egects of the stored energy is still effective because there was much less
intralath 8 in the unrecrystallized-transforrned M steel than in the recrystallized M steel.
Td= 1080 O C -- < 5 % aBL -- 8-bearing -- 100%aBL -- y grain boundaries are clex -- little intragranular nucleation -- intralath 8
In CCT M steel, there is no obvious change in a~ morphologies and CCT diagrarns
because both GBa and ae transformation are promoted by deformation below TNR. The
7'' = 850 'C - - 100 % aeA -- almost 8-fiee - - 100 % aeL -- subgrain boundaries are clear -- deformation bands are seen -- tittle intraiath 8
a~ laths are thin and almost fiee of iatraiath û because of the same dislocation pipe
diffision as in L steel.
7.5 BAINITE TRANSFORMATION MODELS
Based on the microstructural observations, a mode1 accounting for a~ nucleation
and growth is proposed for each steel.
7.5.1 L Steel As-Reheated
The mode1 for ae transformation in L steel as-reheated is iilustrated in Figures
7.5.1 through 7.5.4.
7.5.1 - 1 Nucleation
ae occurs over a wide cooling rate range, 130-1 "Ch. At intermediate cooling
rates, ae nuclei form at y grah boudaries (Figure 7.5. la). At high cooling rate, as also
aucleates at twin boundaries due to large undercooling (Figure 7.5. lb). At slow cooling
rates, a~ nuclei form at y grain boundaries and, intragranularly, at dislocations or C-
depleted areas fonned due to composition fluctuations (Figure 7.5. l c )
To diminish the energy barrier to nucleation, a~ usually keeps a K-S or N-W
orientation relationship with y.
7.5.1 -2 Lenahening
Upon continuous cooling in the range 130-10 OC/s, ae nuclei lengthen into a a L by
a displacive mechanism with some C trapped in a~ laths (Figure 7.5.2a). The
syrnpatheticaily nucleated ae nuclei grow into bundles until irnpingement at grain
boundaries, twin boundaries and pre-existing a~ occurs (Figure 7.5.2b). Within an a~
bundle, aeL laths are separated by low-angle boundaries.
7.5.1.3 Thickening
During and d e r a~ lengthening, C partitioning and 8 precipitation competes
(Figure 7.5.3a) and a~ thickening occurs by the diffisional ledge movement (Figure
7.5.3b) andor the coaiescence of laths through a low-angle boundary migration (Figure
7.5.3~). The former should be dominant at higher temperatures as the diffision of
substitutional atoms are required, where the latter should be controliing at lower
temperatures when C diffusion is still high.
7.5.1.4 Formation of Bainite Morpholoeies
At high temperatures, C-partitioning outweighs 0-precipitation and C difises
rapidly Born C - e ~ c h e d as laths into surrounding y at higher temperatures, leading to Bi.
Its width is increased rapidly by ledge movement at the relatively higher temperature . At
room temperature, some C-enriched y remains, but rnost transforms into au (Figure
7.5.4a). Rarely when C concentration in y exceeds the extrapolated y/B composition
boundary, interlath 0 precipitates to form Bn (Figure 7.5.4b). Upon further cooling, a~
lengthening accelerates due to increased undercooling, intralath 8 precipitation is
dominant and a~ is present mostly as Bm (Figure 7.5.4~). The explanations are: 1) C
diffusivity decreases exponentially with the decreasing temperature, 2) decarburization of
a large arnount of newly-formed a~ laths requires longer range C diffision, and 3) the
remaining y is C - e ~ c h e d due to the formation of BI, which hinders C diffision. Like
t empered a ~ , intralath 0 keeps the Bagary atsky orientation relationship with as.
Thickening of Bm is limited because of lower temperatures, so Bm laths are generafly
thinner than Bi laths. A high density of transformation accommodation dislocations is
introduced following the precipitation of intralath 0, and therefore the intralath 0 coarsens
rapidly.
Lathless aeA fons oniy at slow cooling rates and higher temperatures. Granular
B: first grows via a reconstructive ledge growth mechanism, engulfing y grain boundaries.
Long-range diffusion of C atoms dominants, resulting in w / y ~ islands (Figure 7.5.4d). At
lower temperatures, elongated B: foms by lengthening displacively and thickening
reconstmctively a d o r by sublath coaiescence, resulting in wavy and coarse B: laths that
are slightly misoriented (say, 2 O ) (Figure 7.5 Ae).
7.5.2 L Steel Deformed Above Tm
Compared with that in L steel as-reheated, in L steel deformed above TVR it is
possible to nucleate ae precipitates at twin boundaries even at a lower undercooling.
However, refined y grains result in shortened qBL laths compared with the as-reheated y,
and the depletion of Nb due to precipitation will decrease the as growth rate, leading to a
decrease in CCT B,.
bainite nucleus autenite grain boundary
Figure 7.5.1 ~ucléation of bauiite in L steel as-reheated at: a) intermediate cooling rate, b) hi& cooling rate, c) slow cooling rate.
Figure 7.5.2 Lengthening of bainite in L steel as-reheated.
ledge growth direction
Figure 7.5.3 Thickening of bainite Iaths. a) C partitioning and 8 precipitation, b) ledge growth, c) lath coalescence.
carbide
dtb J 4f #>
d e r carbide precipitaion
carbide
b) B,
\ &er carbide coarsening
te grain bound
retahed austenitehartensite
granular f& te - elongated ferrite
Figure 7.5.4 Formation of bainite morphoiogies. a) BI, b) BD, c) Bm and its 8 coarsening, d) BP, e) BI!
7.5.3 L Steel Deformed Below TNR
At high cooling rates where dislocation subgrains can not fom, a~ nucleates
m d y at y grain boundaries, twin boundaries, and deformation bands (Figures 7.5.5a and
b). At intermediate cooling rates, cellular dislocation subgraios form. Similar to y grain
boundaries, subgrain boundaries serve as nucleation sites of a~ (Figure 7.5.5~). Such
added nucleation sites plus distorted boundaries increase the rate of nucleation and growth
significantiy, resulting in an acceleration of a~ transformations. An UB bundle usually
grows rapidly within a subgrain. This is why there is normally one as bundle in a subgrain
(Figure 7.5.5d).
There are a high density of defonnation-iotroduced dislocations and vacancies,
thereby long-range C difision is enhanced sigaificantly due to dislocation pipe diffusion.
Therefore, compared with recrystallized L steei, unrecrystallized L steel has more a&
constituents and wider a~ laths, mostly BB? (850 OC-deformation) or al1 aeA (780 O C -
deformation), and sirnultaneously, Bm is decreased or eliminated.
7.5.4 M Steel As-Reheated
A mode1 for ae transformations in M steel as-reheated is illustrated in Figures 7.5.6
through 7-53.
Subgrain is not yet formed &gh cooling rate)
bainite nucleus
defornation band grain boun*
a)
Subgrain is completed (intermediate cooling rate)
Figure 7.5.5 Nucleation of bainite in L steel deformed below TM. a) high cooling rates, b) completion of as growth in a), c) intemediate cooling rates, d) completion of ae growth in c).
7.5.4.1 Nucleation
In M steel, a~ occurs only in a narrow intermediate cooling rates (Le., 10-1 "Ch).
A layer of GBa is always present, keeping the K-S orientation relationship with y. It is
energetically favourable for a~ nuclei to form at GBa, still keeping a K-S relationship
with y via GBa. Further cooling results in the formation of aa nuclei at intragranular sites
such as large precipitates and anneaiing twin boundaries (Figure 7-56}.
7.5.4.2 Lengthening
Upon continuous cooling, GBa-formed a w lengthens into ae laths first (Figure
7.5.7), and intragranularly nucleated a~ nuclei lengthen displacively into the so-called
"rnidribs" of Bw laths (Figure 7.5.8a). C atoms diffise from Brv rnidribs into the
surrounding bulk y rapidly (Figure 7.5.8b). Frorn the broad faces of the decarburized
rnidribs, an array of secondary ae can grow, lying - 55-60° to the lengthening direction of
the midrib (Figure 7.5.8~). C then diffuses into gaps of y between these secondary ae
laths and precipitates out as 8. The thickening of the secondary ae laths encloses the
interlath 0 as "intralath" 0. A BN is thus formed into a lenticular morphology (Figure
7.5.8d). New ae nuclei then form and lengthen fiom pre-existing a~ laths.
7.5.5 M Steel Deformed Above ir,,
In this case, more intragranular nucleation sites are produced such as twin
boundaries and strain-induced precipitates, but ae lengthening rate is decreased.
7.5.6 M Steel Deformed Below Tm
It is noted that the intralath 8 is decreased.
grain bounciary fenite
a)NucIeation
c) New bainite formation
b) Lengthening
Figure 7.5.6 Nucleation of bainite in M steel as-reheated
C diffusion \ bainite/Widmanstatten f&te
Figure 7.5.7 Lengthening of GBa-nucleated bainite in M steel as-reheated.
martensitdretained aus teni te
martensitdretained austenite
martensitelretained austenite
4
Figure 7.5.8 Formation of Bw in M steel as-reheated. a) midrib formation, b) C atoms diffusion and 8 precipitation, C) secondary as laths growth, d) BE completion.
CONCLUSIONS
in this investigation, the effects of austenite condition on bainite transformations
were studied for a low-C @,Nb)-rnicrodloyed bar steel and a medium-C Nb-microailoyed
forging steel. Various austenite conditions were obtained mainly by changing the
deformation temperature above and below Tm. Each austenite condition was
characterized as t O grain size, precipitate distribution and sub structure. B ainite
transformation kinetics were determined by dilatometry for continuous cooiing and
isothermal transformations. Bainite microstmctures were anaiyzed using TEM and SEM,
and sorne bainite transformation models were proposed. The following conclusions are
made:
Austenite Condition
1) Most of the Nb and B is in solution at the 1180 O C reheat temperature, and
austenite defomation produces strain-induced precipitates in both steels. These
precipitates are generally Nb-nch. Precipitates forming at y grain bounduies and on the
deformation substmcture (defomation bands, subgrains) are coarser than the matrix
precipitates due to faster diffusion at grain boundaries and deforrnation substructures. In
addition, precipitates seem to form in clusters following defomation above Tm.
Bainite TES
Distinct bainite types BE BI, Br, Bn, Bm, and BIV represent decreasing
uansformation temperature, decreasing carbon-diffusion distance, and a transition from
reconstmctive to displacive transformation mechansim.
Bainite Transformation Mechanisms
1) In recrystallized austenite, bainitic femte nucleates at grain boundaries or grain
boundary ferrite, and intragranularly at twin boundaries, precipitates and pre-existing
bainitic femte Iaths.
2) In deformed-unrecrystailized austenite, bainitic ferrite nucleates at deformed
austenite grain boundaries, grain boundary femte, and intragranularly at subgrain
boundaries, noncoherent twin boundaries, deformation bands, strain-induced precipitates
and pre-existing bainitic femte laths.
3) Lathlike bainite lengthens displacively and thickens by a difision-controlled
ledge mechanism andfor by subiath coalescence.
4) Lathiess bainite grows by either a displacive or reconstmctive mechanism.
Effects of Austenite Condition
1) For deformation above TNR, bainite transformation is slightly retarded due to
delayed bainite growth caused by strain-induced precipitation.
2) For deformaiion below Tm bainite transformation is significantly accelerated
due to increased austenite grain surfaces and deformation-introduced subgrain boundaries,
deformation bands, and noncoherent twin boundaries.
3 ) In the case of deformation below Tm, Iathless bainite replaces lathlike bainite
largely, leading to raised B,. Carbides are alrnost or completely eliminated due to the
significantly increased carbon partitionhg by dislocation pipe dinusion in the investigated
cooling rate range.
SUGGESTIONS FOR FUTURE WORK
Austenite Conditions
This study used oniy one reheat (austenitization) temperature of 1180 O C . A
limited study using 900 OC and 1250 OC reheat temperatures demonstrated a significant
change in bainite microstructures. For example, L steel reheated at 1250 O C contained a
large volume fraction of intragranular bainite la th (maidy B3, while L steel reheated at
900 O C generally had a lathless bainite morphology regardless of the cooling rate.
Therefore a detailed study of a wider range of reheat temperatures and austenite
conditions is suggested.
Bainite Transformation Mechanisrns
1. It has been observed fiequently that bainite laths are ofien distributed among
several variants, and bainite microstructures are aiigned after deformation below Tm.
These variants can be identified by the trace analysis method [Sandvik, 198 la] or by the
electron backscattering pattern (EBSP) technique to provide information on the
crystdlography of bainite transformations. M steel is a good materiai for such a study as
individual Brv laths are distributed on a martensite rnatrix.
2. The nucleation mechanisms of granula BP that forms at high temperatures are not
yet clear. A further study in this regard is suggested.
3. Boron is believed to have significantly influenced the formation of both ferrite and
bainite. A study of the distribution of boron would be useful.
4. Such interphase interfaces as bainitic femte Iathdretained austenite, bainitic femte
lathdintrdath carbides, and bainitic ferrite lathdprecipitates could be revealed using high
resolution TEM. To preserve enough austenite at room temperature, an isothemally
transformed sample instead of a continuously transformed one is recommended.
5 . This study has coiiected a vast amount of data on bainite transformation kinetics.
At this stage, a kinetic mode1 for bainite transformation could be developed, which relates
austenite condition to the desired bainite microstmcture.
Mechanical Properties
1. There are a nurnber of different lathless bainite microstructures produced by L
steel as-reheated and slowly cooled, L steel deformed below Tm and intermediately
cooled, or L steel reheated at Iower temperatures (Le., 900 OC) regardless of cooling rate.
These microstnicnires are believed to have attractive mechanicd properties, and should be
studied.
2. In the case of deformation below Tm the formation of subgrains (at an
intermediate cooling rate for L steel or isothermal holding for M steel) effectvely refines
austenite grains and produces a bainite microstructure featuring few or no carbides. This
may provide a good combination of strength uid toughness, and a study of the mechanicd
propenies for such microstructures is recommended.
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CALCULATION OF EQUILIBrWM VOLUME FRACIlON OF SIXAIN-
INDUCED Nb(C,N) PRECIPH'ATE3
For Fe-Nb-C-N quaternary system [Rios, 19881:
where XN~, xs and x~ are mole fractions of solute Nb, C and N in y, respectively. 'yc and
)N are the site fiactions of C and N in Nb(CN), respectively, and y= + y~ = 1 when the
Nb(C,N) phase is expressed by the two-sublattice mode1 mllert and Staffmson, 19701.
k N ~ and km are the solubility products of M C and NbN, respectively. From Irvine et al.
[ 1 9671 and Balasubr&an and Kirkaldy [ 1 9891 :
By mass balance:
where >YN6, 'xc and are the initial mole fiactions of Nb, C and N.
Therefore, the eqiillibrium volume &action V) of strain-induced Nb(C,N) precipitates
can be obtained by solving Equations Al through A4.
The driving force for nucleation, AG, can be caiculated according to the equation
provided by Liu and Jonas 11 9891:
where R is the gas constant (8.3 14 J/mol.K), and T the reheat temperature O().
Ifyc = y ~ , Equation AS is converted into the simpler Equation 2.6.
